Rebecca A. MacKay

Glenn Research Center, Cleveland, Ohio

Stephen W. Smith

Langley Research Center, Hampton, Virginia

Sandeep R. Shah

George C. Marshall Space Flight Center, Marshall Space Flight Center, Alabama

Robert S. Piascik

NASA Engineering and Safety Center, Langley Research Center, Hampton, Virginia

Reaction Control System Thruster Cracking

Consultation: NASA Engineering and Safety

Center (NESC) Materials Super Problem

Resolution Team (SPRT) Findings

NASA/TP—2005-214053

December 2005

Rebecca A. MacKay

Glenn Research Center, Cleveland, Ohio

Stephen W. Smith

Langley Research Center, Hampton, Virginia

Sandeep R. Shah

George C. Marshall Space Flight Center, Marshall Space Flight Center, Alabama

Robert S. Piascik

NASA Engineering and Safety Center, Langley Research Center, Hampton, Virginia

Reaction Control System Thruster Cracking

Consultation: NASA Engineering and Safety

Center (NESC) Materials Super Problem

Resolution Team (SPRT) Findings

NASA/TP—2005-214053

December 2005

National Aeronautics and

Space Administration

Glenn Research Center
Available from

NASA Center for Aerospace Information

7121 Standard Drive

Hanover, MD 21076

National Technical Information Service

5285 Port Royal Road

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Trade names or manufacturers’ names are used in this report for

identification only. This usage does not constitute an official

endorsement, either expressed or implied, by the National

Aeronautics and Space Administration.

Available electronically at http://gltrs.grc.nasa.gov

Acknowledgments

This paper is the result of considerable effort from a number of dedicated NASA civil servants and support service

contractors. The first author would like to thank her coauthors, particularly for their hard work and continuous

support throughout this investigation, as well as for their insight into environmental cracking mechanisms.

The first author also wishes to acknowledge Mr. Robert H. Titran (retired, Glenn Research Center (GRC)) for

sharing his 30 years of refractory metals experience and for being the first to suggest that the thruster cracking was

more likely the result of hydrogen embrittlement; Mr. David R. Hull (GRC), Mr. James W. Smith (QSS, GRC), and

Mr. Terry R. McCue (QSS, GRC) for their dedication and for their considerable expertise provided during the

failure analyses of S/Ns 120 and 132; Dr. Hugh R. Gray and Dr. Michael V. Nathal (GRC) for encouragement and

helpful suggestions; Dr. Binayak Panda and Mr. Gregory A. Jerman (Marshall Space Flight Center) for their

expertise and technical support during the failure analyses; and Mr. Terence Kelly (Boeing White Sands),

Mr. Michael Tarkanian (Boeing Huntington Beach), and Dr. Sherwin Yang (Boeing Rocketdyne) for their

cooperation and technical support.
NASA/TP—2005-214053 1 

Reaction Control System Thruster Cracking Consultation: 

NASA Engineering and Safety Center (NESC) Materials 

Super Problem Resolution Team (SPRT) Findings 

Rebecca A. MacKay 

National Aeronautics and Space Administration 

Glenn Research Center 

Cleveland, Ohio 44135 

Stephen W. Smith 

National Aeronautics and Space Administration 

Langley Research Center 

Hampton, Virginia 23681 

Sandeep R. Shah 

National Aeronautics and Space Administration 

George C. Marshall Space Flight Center 

Marshall Space Flight Center, Alabama 35812 

Robert S. Piascik 

NASA Engineering and Safety Center (NESC) 

Langley Research Center 

Hampton, Virginia 23681 

Summary 

The shuttle orbiter’s reaction control system (RCS) primary 

thruster serial number (S/N) 120 was found to contain cracks 

in the counter bores and relief radius after a chamber repair 

and rejuvenation was performed in April 2004. Relief radius 

cracking had been observed in the 1970s and 1980s in seven 

thrusters prior to flight; however, counter bore cracking had 

never been seen previously in RCS thrusters. Members of the 

Materials Super Problem Resolution Team (SPRT) of the 

NASA Engineering and Safety Center (NESC) conducted a 

detailed review of the relevant literature and of the documentation from the previous RCS thruster failure analyses of the 

1980s. It was concluded from the review that the previous 

failure analyses lacked sufficient documentation to support the 

conclusions that stress corrosion cracking or hot-salt cracking 

was the root cause of the thruster cracking and lacked reliable 

inspection controls to prevent cracked thrusters from entering 

the fleet. The NESC team identified new materials characterization and mechanical tests focused specifically on determining (a) the environmental cracking damage mode (root cause) 

in the niobium RCS thrusters and (b) the likelihood of growth 

of environmental cracking while in service (after manufacturing). To this end, the NESC team, comprising the NASA 

Glenn Research Center (GRC), the NASA Langley Research 

Center (LaRC), and the Marshall Space Flight Center (MSFC) 

researchers, performed detailed materials characterization on 

thruster hardware (S/Ns 120 and 132), as well as pertinent 

testing and characterization of niobium alloy specimens. The 

NESC team also pursued technical interchanges and conducted 

testing and evaluations with Safety and Mission Assurance 

and the Orbiter Materials and Processing group to address concerns with thruster repair and rejuvenation processes at the 

White Sands Test Facility (WSTF) that have the potential of 

embrittling the thruster material under certain conditions. 

The purpose of this report is to summarize the major findings and to provide the supporting documentation for the RCS 

thruster cracking consultation by the NESC team. The following findings will be shown and substantiated: 

 (1) RCS thruster intergranular cracking (IGC) is due to 

hydrogen embrittlement. 

 (2) Thruster IGC was produced during manufacturing and 

was the result of processing the thrusters in the presence of 

fluoride-containing acids. 

 (3) Appreciable environmental crack propagation does not 

occur after manufacturing. 

 (4) The thruster material was not significantly affected by 

the sodium hydroxide immersion during the repair and rejuvenation processes at WSTF. 

 (5) A recommendation was made to audit the thruster 

repair and rejuvenation processes at WSTF to eliminate steps 

that have the potential of allowing additional hydrogen uptake 

and inducing hydrogen embrittlement. 
NASA/TP—2005-214053 2 

Section 1: Introduction and Background 

to Thruster Cracking 

Thirty-eight primary reaction control system (RCS) thrusters are located in forward and aft positions of each shuttle 

orbiter. These thrusters are used to provide attitude (pitch, 

yaw, and roll) maneuvers, as well as translation maneuvers 

along the orbiter axis. Thruster serial number (S/N) 120 was 

undergoing a chamber repair and rejuvenation in early 2004 at 

White Sands Test Facility (WSTF) as a result of coating spallation. After the chamber repair, the chamber and nozzle 

assembly was electron beam welded to the old injector. During the post-weld, dye-penetrant inspection of the new chamber-to-injector weld, cracks were found in the counter bores 

and subsequently in the relief radius of the niobium (formerly 

known as columbium (Cb)) alloy C−103 injector. No cracks 

had ever been detected before in the counter bore locations, 

although cracks were detected between 1979 and 1982 in the 

relief radius of seven injectors (refs. 1 and 2). Figure 1 (ref. 3) 

is a schematic diagram of a longitudinal cross section of the 

injector; the locations of the relief radius and counter bores are 

indicated on opposite sides of the niobium injector flange. The 

thruster cracking is shown by the dye penetrant indications in 

figures 2 and 3 (ref. 4) for the counter bores and relief radius, 

respectively, in S/N 120. It is evident from figure 3 that a large 

fraction of the circumference of the relief radius was cracked 

in S/N 120. The initial concern was that hot combustion gases 

could leak to the outside of the thruster and affect the shuttle 

orbiter in a Criticality 1/1 (i.e., loss of vehicle with no redundancies) failure mode (ref. 3) if the cracks had propagated 

entirely through the injector wall to the acoustic cavities. 

The prior failure analyses from 1982 concluded that the 

cause of the relief radius cracking in the RCS thrusters was 

stress corrosion, as a result of the use of a preweld etchant that 

contained hydrofluoric acid (HF) and the fit-up stresses from 

torquing the bolts that secure out-of-flat flanges (ref. 1). The 

excess etchant was believed to have run through the bolt holes 

and to have become entrapped in the relief radius, since the 

mounting flange could not be lifted to adequately rinse the 

relief radius after the etchant was used. The prior failure 

analyses also report that the cracks were intergranular all the 

way to the crack tip, that CbF3 was present on the crack surface, and that all cracks initiated during manufacturing, specifically during the 316 °C (600 °F)/48-hr air bakeout of the 

chamber insulation (ref. 1). In this prior investigation, the 

CbF3 reaction product was incorrectly identified, and should 

have been identified as NbF5 (refs. 5 and 6). The prior failure 

analyses testing reproduced the IGC in C−103 when the 

preweld etchant was dried on a niobium test specimen; the 

specimen was then covered with a titanium plate to entrap the 

etchant, stressed in tension, and heated to 316 °C 

(600 °F)/48 hr to simulate the elevated temperature bakeout 

manufacturing process. It should be noted that the thruster 

failure was described later as “hot fluoride salt stress corrosion 

cracking” in a summary article on the RCS thruster failure 

(ref. 7), although the term “hot salt” was not used to describe 

the thruster failures in the original failure analysis 

documentation. 

S/N 120 was submitted in its entirety to the M&P organization at Boeing Huntington Beach (HB) for failure analysis in 

June 2004. The portion of S/N 120 that was set aside for 

destructive analysis is shown in figure 4 (ref. 4); the remainder 

of the injector was left intact with the combustion chamber 

and was used in nondestructive evaluation (NDE) development efforts. The Boeing HB failure analysis (ref. 8) concluded that the injector cracking was intergranular and that the 

overall appearance of the cracking in S/N 120 was similar to 

what was described in the failure analyses conducted in the 

early 1980s. The presence of carbon, oxygen, and fluorine was 

found on the fracture surfaces, although the age of the cracks 

could not be determined. The chemical analysis, grain size, 

and hardness of the bulk material (refs. 4 and 8) met the original material specifications, which are given in table I (ref. 9). 

In addition, no evidence of overload by ductile rupture or of 

fatigue rupture was found (ref. 4). 

The detailed review of the relevant literature and of the 

documentation from the previous RCS thruster failure analyses of the 1970s and 1980s provided little evidence in support 

of stress corrosion cracking (SCC) or hot-salt cracking in 

niobium alloys. Furthermore, it was concluded that the previous failure analyses lacked sufficient documentation to support the conclusions that SCC (ref. 1) or hot-salt cracking 

(ref. 7) was the root cause of the thruster cracking. Since a 

large fraction of the thruster fleet was manufactured during the 

time when the HF-containing, preweld etchant was used, a 

large fraction of the thruster fleet is presumed to have cracking 

damage already present. No reliable NDE technique is currently available for detecting relief radius or counter bore 

cracking during an in situ examination of the thrusters 

installed on orbiter vehicles. Thus, the NASA Engineering and 

Safety Center (NESC) team questioned if the environmental 

damage mode had been truly identified in the RCS thrusters 

and if the thrusters could operate safely with environmental 

cracking damage modes. Specific recommendations were 

made by the NESC team, which included additional materials 

characterization and mechanical testing to determine (1) the 

root cause of the thruster cracking and (2) the likelihood of 

growth of existing thruster cracks while “in service” (after 

manufacturing). It should be noted that the Shuttle Program 

Office agreed with the NESC team recommendations and 

funded the testing and characterization that supported this 

work. Researchers at Glenn Research Center (GRC), Langley 

Research Center (LaRC), and Marshall Space Flight Center 

(MSFC) performed detailed materials characterization on the 

thruster hardware and mechanical testing on the niobium 

C−103 alloy to reproduce the intergranular failure and to 
NASA/TP—2005-214053 3 

determine if crack propagation was likely. The purpose of this 

report is to present the major findings and supporting documentation for the RCS thruster cracking consultation by the 

NESC Materials SPRT. 

Section 2: Root Cause Analyses and 

Testing 

2.1 Literature Review 

The prior failure analyses from 1982 concluded that the 

thruster cracking was caused by stress corrosion (ref. 1); this 

was later described as hot-salt cracking (ref. 7). A detailed 

literature review was conducted that showed little evidence for 

SCC or hot-salt cracking in niobium, and in fact, there was an 

absence of literature on SCC and hot-salt cracking in niobium. 

The only literature found in an exhaustive search that described SCC in niobium was two papers that described the 

RCS thruster case from the 1980s (refs. 7 and 10). Discussions 

with refractory metals experts (refs. 11 to 13) and the review 

of the literature show that niobium was well known to be 

susceptible to hydrogen embrittlement (refs. 14 to 28), as well 

as to interstitial contamination from oxygen and nitrogen. This 

result provided the impetus for electrochemical testing (see 

section 2.4) to verify that hydrogen production at niobium 

surfaces was thermodynamically viable and to determine the 

environmental conditions under which thruster cracking would 

likely occur. 

Both intergranular stress corrosion and hot-salt cracking 

require a susceptible grain boundary chemistry, which has not 

been observed in the literature for niobium. Niobium alloys 

generally have very homogeneous microstructures and therefore very little variation in chemistry near the grain boundaries 

would be available to drive a localized corrosion process via 

SCC or hot-salt cracking. Specifically, grain boundary segregation would need to be present in the form of either grain 

boundary precipitates or solutes at the grain boundaries for 

preferential anodic dissolution of material at the grain boundaries to occur (refs. 28 to 31). If segregation at the grain 

boundaries is not present in a material, SCC or hot-salt cracking would be highly unlikely. Grain boundary segregation can 

be readily detected by electron microprobe analysis. Therefore, based on the literature review, electron microprobe 

analysis was required to determine root cause and was incorporated into the characterization plan. 

2.2 Electron Microprobe Analyses 

Electron microprobe dot maps, line scans, and point analyses were generated to determine the location of specific elements within the intergranular cracks and within the 

uncracked microstructure. Dot maps enabled broad areas of 

the microstructure to be analyzed, so that general trends in 

composition could first be established for numerous locations 

within the thruster material. Line scans are generally more 

sensitive to small chemistry variations than dot maps and were 

used to confirm dot map data across cracks and grain boundaries. Point analyses provide the most quantitative data of these 

three microprobe techniques and were performed to determine 

if grain boundary segregation was present, once the general 

trends were clearly understood. 

The following sections of S/N 120 were analyzed by 

microprobe in order to substantiate root cause of thruster failure: (a) relief radius crack near bolt hole J and uncracked 

region adjacent to cracking; (b) counter bore and relief radius 

cracks of bolt hole H and uncracked region adjacent to cracking; (c) small counter bore cracks of bolt hole K and 

uncracked region adjacent to cracking; and (d) small counter 

bore cracks of bolt hole M and uncracked region adjacent to 

cracking. These bolt holes are identified in the injector section 

that was set aside for failure analysis (fig. 4). 

2.2.1 Relief Radius Crack Near Hole J 

A large relief radius crack near bolt hole J was selected for 

initial microprobe examination, and the specific areas (A to E) 

examined along the crack and away from the crack are indicated in figure 5. The full depth of the intergranular crack in 

this location is not known because the tip of the major crack 

was not contained within the cut section provided by Boeing 

HB. Area dot maps were obtained at magnifications of 200×

and 1000× near the relief radius surface (area A), as well as 

~1.8 mm (0.07 in.) away (areas B and C) from the surface, in 

an effort to determine if any qualitative differences in the 

distributions of elements were present as a function of distance 

from the relief radius surface. Uncracked areas (near D and E) 

were also examined to detect the presence of precipitates or 

higher concentrations of solutes, if any, along grain boundaries away from cracks. Such features, if present, would indicate a predisposition to SCC or hot-salt cracking. 

Figure 6(a) shows a backscattered electron (BSE) image of 

the crack near the surface of the relief radius. BSE imaging 

clearly delineates the cracks and reveals atomic number differences within the microstructure. The major cracks follow 

the grain boundaries in the material; finer cracks are also 

observed to run nearly parallel to the relief radius surface 

along the grain boundaries beneath the deformed layer from 

machining. Figures 6(b) to (d) show the dot maps for niobium, 

hafnium, and titanium, respectively, for the same field of 

view. The qualitative concentration levels are given for each 

element in the keys to the right of the figure; the brighter areas 

in the dot maps indicate more elevated levels of each respective element. Careful examination of figure 6(b) shows that 

the niobium levels lining the surfaces of the cracks are 

reduced from the levels within the general microstructure; this 

result is from the niobium being tied up in the oxide on the 

crack surfaces. Hafnium and titanium are not evident within 
NASA/TP—2005-214053 4 

the cracks (figs. 6(c) to (d)). Figures 6(b) to (d) also show that 

the grain boundaries in uncracked areas are not delineated by 

the major elements (niobium, hafnium, and titanium). The 

most prominent feature in figures 6(a) to (d) is the macroscopic banding that occurs for these major elements. It is 

apparent that slightly elevated concentrations of hafnium and 

titanium reside in bands wherever the niobium is somewhat 

leaner, and vice versa. This was consistently observed in other 

areas of the injector and represents the residual forged 

macrostructure. 

Figure 7 represents the same near-surface area as that presented in figure 6 and contains dot maps for fluorine, carbon, 

oxygen, and nitrogen. These elements were examined to confirm the presence of elevated levels of fluorine, carbon, and 

oxygen that were found on the crack surfaces in the prior 

failure analysis (refs. 4 and 8) and also to examine for the 

presence of these elements along grain boundaries in uncracked regions, since oxygen and nitrogen are known to 

cause embrittlement in niobium materials (refs. 11, 12, 14, 16, 

18, and 24). Hydrogen cannot be detected by the electron 

microprobe. Figure 7(a) shows that fluorine is present along 

the cracked intergranular surfaces and along part of the relief 

radius surface, as indicated by the bright blue-green levels. 

Carbon and oxygen are more readily discernable within the 

crack (figs. 7(b) and (c)). Some of these levels of carbon and 

oxygen may have been enhanced by the limited infiltration of 

the mounting medium from the relief radius surface, although 

carbon and oxygen are also elevated near the crack tip, where 

infiltration of the mounting medium did not occur. Thus, the 

elevated levels of carbon and oxygen on the intergranular 

fracture surfaces in these areas of S/N 120 are consistent with 

previously reported scanning Auger results (refs. 4 and 8). 

Elevated levels of nitrogen were typically not apparent along 

the cracks (fig. 7(d)). It is important to note that the microprobe dot maps do not indicate a presence of fluorine, carbon, 

oxygen, and nitrogen outside the cracks, since no elevated 

levels of these elements were observed outside the cracks that 

were above background levels. Thus, no local segregation of 

fluorine, carbon, oxygen, and nitrogen was present outside the 

cracks. Dot maps were also obtained for zirconium, tantalum, 

and tungsten; these elements were not found within the cracks 

and were not found to be segregated within the uncracked 

microstructure. 

Similar results were obtained from regions further in from 

the machined surfaces along the main crack and at terminations of side branches (figs. 8 to 11). Again, niobium levels 

within the cracks were reduced when compared to the general 

microstructure, which is the result of niobium oxide on the 

crack surfaces. Hafnium and titanium were again not found 

within the cracks. Fluorine, carbon, and oxygen were clearly 

present along the intergranular cracks to the crack terminations, with oxygen being observed as a continuous film along 

the entire crack surface. Fluorine and carbon were not 

consistently observed as a continuous film along the intergranular cracks. 

Microprobe line scans were also conducted across relief 

radius cracks. Line scans provide counts as a function of distance for each element selected and are more sensitive to small 

variations in chemistry than microprobe dot maps. The line 

scans were performed normal to the intergranular cracks in 

polished cross sections so that data could be obtained from 

within the crack itself and from material just outside of the 

crack. An example of the position of a typical line scan is 

indicated by the white arrow superimposed in figure 8(a), and 

figure 12 provides an example of typical data obtained from 

such a line scan. Figure 12(a) provides gray scale image 

counts across the crack that was analyzed; the crack width is 

indicated in this figure where the gray scale counts drop off 

precipitously to zero. Figures 12(b) and (c) represent the analyzed data for fluorine and oxygen, respectively, as a function 

of distance across the intergranular crack. It may be seen that 

the number of counts for each element increases significantly 

within the crack relative to background levels. These line 

scans confirm the dot map data that indicate that higher concentrations of fluorine and oxygen were present within the 

cracks and that these elements did not extend beyond the 

cracks into the parent material. 

Microprobe dot maps were also obtained in regions far 

removed from the intergranular cracks to examine the material 

for the presence of any precipitates or higher concentrations of 

solutes. Features such as these would indicate that the C−103 

may be predisposed to SCC or hot-salt cracking. An example 

of an uncracked region away from the relief radius crack is 

shown in figure 13. The BSE image in figure 13(a) delineates 

several grains in the field of view at a magnification of 1000×. 

No precipitates or elevated levels of the major elements are 

evident on the grain boundaries, although a few hafnium 

oxides are observed within some of the grains. The banded 

macrostructure is visible, particularly in figures 13(c) to (d), 

and the hafnium oxides are observed within a hafnium-rich 

band. Dot maps of fluorine, carbon, oxygen, and nitrogen in 

figure 14 also reveal no elevated concentrations of these elements above background levels. These results were confirmed 

in other uncracked regions of S/N 120 near bolt holes H, K, 

and M. 

2.2.2 Hole H Relief Radius 

The relief radius in the vicinity of bolt hole H contained 

several shallow cracks, which were only up to several grains 

deep in from the surface (fig. 15). These cracks were located 

near the end of the large relief radius crack indication that was 

revealed by the dye penetrant inspection in figure 3 for S/N 

120. A box superimposed on the image in figure 15 highlights 

a crack along the injector flange; microprobe analyses were 

performed on this area to determine if shallow cracks yielded 
NASA/TP—2005-214053 5 

similar results to those obtained on deeper cracks (i.e., near 

hole J). The selected crack in figure 15 completely surrounded 

a grain which subsequently fell out during metallographic 

polishing. The void that was left behind enabled a unique 

opportunity for the analyses of broad surfaces of grain 

boundaries. Figures 16(a) to (d) show the BSE image for the 

analysis area and the dot maps for fluorine, oxygen, and carbon, respectively. It is evident from these data that high concentrations of fluorine, oxygen, and carbon were present on 

the broad grain boundary surfaces of the shallow crack, which 

is consistent with the results from the polished cross sections 

of the deep relief radius crack near bolt hole J. Microprobe 

line scans also confirmed the presence of elevated levels of 

fluorine and oxygen within the shallow intergranular cracks. 

2.2.3 Hole K Counter Bore 

Dye penetrant results from reference 4 indicated that the 

counter bore of hole K was crack free. However, scanning 

electron microscopy revealed the presence of numerous small 

cracks (up to 30 μm (0.001 in.) deep) along the counter bore 

surface of hole K; these were the most shallow cracks observed in S/N 120. These shallow cracks extended through the 

~8 μm (0.0003 in.) thick, cold-worked machining layer at the 

surface. Some cracks proceeded inward along grain boundaries, as in figure 17(a), while other cracks proceeded inward 

along a transgranular path through grain interiors (fig. 17(b)). 

Microprobe dot maps revealed that both the intergranular and 

transgranular cracks were oxidized, as evidenced by the higher 

concentrations of oxygen along the crack surfaces, compared 

to that within the uncracked microstructure (fig. 17(c)). However, based on an analysis of 10 individual cracks, none of 

these cracks were observed to contain fluorine. An example of 

this is shown in figure 17(d). The reason for the consistent 

lack of fluorine in these shallow counter bore cracks is not 

known. It may be speculated that these shallow cracks did not 

result from exposures to HF during manufacturing but instead 

resulted from exposure to another source of hydrogen. A lack 

of fluorine was also exhibited by C−103 specimens believed to 

have been embrittled by hydrogen during C-ring testing (section 2.3). In contrast, the significantly deeper counter bore 

cracks in S/N 120 from holes M (ref. 4) and H, for example, 

did contain significant concentrations of fluorine, carbon, and 

oxygen on the intergranular crack surfaces. So in general, it 

may be stated that the IGC in the counter bores had similar 

features to those in the relief radius, except for the very shallow (~30 μm (0.001 in.)) counter bore cracks. 

 2.2.4. Quantitative Point Analyses 

Uncracked regions of S/N 120 were examined quantitatively for the determination of chemical segregation, if any, to 

the grain boundaries. Electron microprobe data were obtained 

using point analyses in the vicinity of the relief radius and 

counter bore of bolt hole J, the counter bore of bolt hole K, 

and the relief radius and counter bore of bolt hole M. Figure 18 shows a representative area examined in the counter 

bore region of bolt hole K; an overall grain size of ~50 μm 

(0.002 in.) is evident in the field of view. Superimposed on the 

image in figure 18 are the specific locations that were used for 

the point analyses obtained from 20 grain boundaries and 20 

neighboring grain interiors. The means and 95 percent confidence intervals of the levels of the major elements, niobium 

and hafnium, are presented in figures 19(a) and (b), respectively. The means for the grain interiors and the grain boundaries are indicated in the figures by the filled circles and 

unfilled squares, respectively. It is evident in figures 19(a) and 

(b) that the major element levels in the grains and along the 

grain boundaries are nearly identical and that their confidence 

intervals overlap. Thus, the data clearly indicate that the compositions of the major elements are statistically equivalent for 

the grain boundaries and grain interiors. 

Figure 19(c) shows the compositional levels for the minor 

elements and interstitials from the same region of the hole K 

counter bore. The data clearly indicate no difference in the 

carbon, zirconium, titanium, oxygen, and tantalum contents 

between the grain interiors and the grain boundaries. Tungsten 

was not specifically examined in this sample because the element was not found to segregate in other areas of S/N 120. 

Additionally, fluorine and nitrogen were examined but were 

not detected in any of the uncracked regions and thus were not 

included in the plot. The analyzed levels for carbon and oxygen are higher than those obtained from the C−103 bulk 

chemical analyses, because the microprobe standards for carbon and oxygen were not within the range of these interstitials 

in C−103. To obtain more appropriate standards for carbon 

and oxygen would have required the procurement and analysis 

of additional standards, and this was beyond the scope of this 

particular investigation. However, the determination of relative differences in carbon and oxygen between grain boundaries and grain interiors is sufficient for the identification of 

grain boundary segregation. The results demonstrate that there 

is no difference between the chemistries of the grains and 

grain boundaries in the injector material. Additional point 

analyses data from the relief radius and counter bore of bolt 

hole J and the relief radius and counter bore of bolt hole M 

substantiated the results from the counter bore of hole K. 

Thus, the major, minor, and interstitial elements are not segregated within the C−103 thruster material, and therefore, cracking by intergranular stress corrosion or by hot-salt cracking is 

highly unlikely. 

2.2.5 Summary and Conclusions for Microprobe Analyses 

Cracking by intergranular stress corrosion or by hot-salt 

cracking is highly unlikely, because no grain boundary segregation was present in the thruster material. As a result, the 

root cause identified in the prior failure analyses (ref. 1) and 
NASA/TP—2005-214053 6 

subsequent reviews (refs. 7 and 10) of the thruster case study 

cannot be substantiated by the current work. 

(1) The C−103 thruster material was determined to be 

homogeneous, and no grain boundary precipitates were 

observed. 

(2) Quantitative data obtained from over 100 data points 

from relief radius and counter bore regions show that the concentrations of the elements at the uncracked grain boundaries 

were statistically equivalent to those within the grain interiors. 

(3) Electron microprobe data from dot maps, line scans, 

and point analyses indicate that oxygen, carbon, and fluorine 

were present within the intergranular cracks, which is consistent with the Auger results from Boeing HB (refs. 4 and 8). 

(4) No elemental segregation was observed in the parent 

material ahead of the crack tips. 

2.3 Testing Attempts to Reproduce Hot-Salt Cracking 

2.3.1 Testing Procedure 

One of the root causes reported in a review article to be 

responsible for the IGC observed in the niobium injectors was 

termed “hot-salt cracking” (ref. 7). In the fall of 2004, Boeing 

Rocketdyne made an attempt to reproduce hot-salt cracking in 

the laboratory on C−103 test coupons (ref. 32). This approach 

was intended to introduce hot fluoride salts to the niobium 

material while precluding the introduction of hydrogen. C-ring 

test specimens were machined and prepared by dipping their 

preloaded tensile surfaces into molten NbF5 at 93 °C (200 °F). 

Since NbF5 is hygroscopic (ref. 32), the hot-salt application 

was performed in an argon glove box. Samples were cooled to 

ambient temperature and were exposed isothermally in air for 

48 hr at temperatures of 149, 235, and 316 °C (300, 455, and 

600 °F) under stresses ranging from 103 to 207 MPa (15 to 

30 ksi). The 149 and 235 °C (300 and 455 °F) temperatures 

were selected for C-ring testing because they represent typical 

injector temperatures in the vicinity of the cracking, with 

235 °C (455 °F) being the maximum under normal operating 

conditions; 316 °C (600 °F) was selected to simulate the 48-hr 

insulation bakeout that thrusters undergo during manufacturing. Optical and scanning electron microscopy (SEM) was 

performed at Boeing Rocketdyne on the C-ring samples, and 

residue from the C-ring surfaces was collected and analyzed 

by x-ray diffraction. Microscopy revealed shallow IGC only in 

the 316 °C/207 MPa (600 °F/30 ksi) specimen. X-ray diffraction identified the specimen surface residue to be NbO2F, 

which was reported to have resulted from a chemical reaction 

between NbF5, oxygen, and residual moisture in the air 

(ref. 32). Boeing Rocketdyne requested that GRC conduct 

microprobe analyses on their tested specimens. The remainder 

of this section will describe the results generated at GRC. 

2.3.2 Microstructural Examination 

All of the C-ring specimens tested at 149, 235, and 316 °C 

(300, 455, and 600 °F) were metallographically prepared and 

examined by SEM. It was confirmed that IGC was present 

only in the 316 °C/207 MPa (600 °F/30 ksi) specimen, and an 

example of the morphology of the cracks is given in figure 20. 

As may be seen in the figure, the cracking was shallow with 

most cracks observed to extend up to a depth of only about 

two grains (~50 μm (0.002 in.)) in from the surface. Additional cracks are also observed to run parallel to the surface, 

beneath the cold-worked layer from machining. Depth profiling of the IGC was performed by controlled polishing iterations, each of which removed 0.25 mm (0.01 in.) of material. 

Scanning electron microscopy was performed after each of 

three polishing iterations to determine the depth and location 

of cracks. It was found that each iteration removed the previously observed cracks and revealed new cracks in different 

locations, which demonstrated that the intergranular cracks 

were neither deep (in from the surface) nor long (into the 

plane of polish). The one exception to this was a single crack 

that intercepted the saw cut along the specimen edge; this 

crack was considerably deeper (~600 μm (0.024 in.)) than all 

of the other cracks and was considerably longer since it was 

observed after each polishing iteration. It is interesting to note 

that all of the cracks observed were located within 2.4 mm 

(0.094 in.) from the cut ends of the specimen (fig. 21), which 

may suggest that the affected material became cracked under 

the action of the cutting process at Boeing Rocketdyne. Otherwise, some cracks should have initiated along the tensile 

surface in the middle of the specimen. 

Electron microprobe dot maps and point analyses were obtained from polished cross sections of the 316 °C (600 °F) 

C-ring specimen. Figure 22(a) shows a typical, surfaceconnected crack in the 316 °C (600 °F) specimen, and dot 

maps for fluorine, carbon, and oxygen from the same specimen area are presented in figures 22(b) to (d), respectively. 

Again, brighter regions in the dot maps are indicative of 

higher concentrations of that element. Although elevated levels of fluorine reside on the specimen tensile surface from the 

fluoride salt that was originally deposited on it, it is clear that 

neither fluorine nor carbon is observed within the crack. This 

is a very different result from that obtained with the intergranular cracks in S/N 120, where fluorine and carbon were 

readily detected within the intergranular cracks. The lack of 

fluorine and carbon in the C-ring specimen was verified by 

other dot maps that were obtained from deeper cracks in the 

C-ring specimen and from other portions of this same crack 

located further away from the surface. In contrast, elevated 

levels of oxygen are consistently observed within the C-ring 

cracks and along the specimen tensile surface, as seen in 

figure 22(d). 
NASA/TP—2005-214053 7 

To confirm the comparative levels of carbon, fluorine, and 

oxygen within the C-ring test specimens, a point analysis 

technique was also performed by microprobe. This technique 

involves the accumulation of counts for a specific element 

over a 10-sec duration at each location examined. Count rates 

obtained from the 316 °C (600 °F) C-ring specimen are displayed in figure 23 for several locations within a crack near a 

triple point, in the C−103 adjacent to the edge of the crack, 

and away from the crack within the C−103 parent material. 

Background count rates at the carbon, fluorine, and oxygen 

peak positions were also obtained within the C−103 parent 

material for comparison. The count rate data confirm the dot 

map results; carbon and fluorine were not evident within the 

C-ring cracks, although elevated levels of oxygen were again 

found within the C-ring cracks. This is unlike the thruster 

cracks that were shown in sections 2.2.1 and 2.2.2 to contain 

elevated levels of carbon, fluorine, and oxygen. To demonstrate that this point analysis technique was capable of detecting carbon and fluorine, if present, count rates were also 

collected for an intergranular crack from S/N 120. These data 

are presented for comparison in figure 24; it is clearly evident 

that elevated levels of carbon, fluorine, and oxygen were able 

to be detected by this technique within the S/N 120 crack. 

Hence, based on polished cross sections of the C-ring cracks, 

fluorine was not evident within the intergranular cracks, even 

though fluoride salts were applied directly to the C−103 tensile surface. 

Although most intergranular cracks were very shallow in 

the 316 °C (600 °F) C-ring specimen, a sizeable crack 

(~600 μm (0.024 in.) deep) was intercepted along the specimen cut end. This crack was deep enough to fracture open in 

order to view any products present on the intergranular surface. The opened crack is shown in figure 25(a), and a higher 

magnification of the grains is given in figure 25(b). The higher 

magnification shows a relatively clean-looking fracture surface, which appears to have no reaction products present. A 

typical Energy Dispersive Spectroscopy (EDS) scan obtained 

from the C-ring fracture surface is displayed in figure 25(c); 

the scan shows high niobium peaks and a sizable hafnium 

peak, which represent the major elements in C−103. The oxygen peak observed is small and barely discernable above 

background levels, which is typical of baseline C−103. No 

fluorine peak is evident, and this result was confirmed in 

numerous locations. For comparison, figures 26(a) to (c) display intergranular crack surfaces, as well as a representative 

EDS scan, obtained from S/N 120. It is clearly evident in 

figure 26(b) that the S/N 120 fracture surfaces contain numerous reaction products, including nodules that resemble opened 

popcorn kernels, as well as layers of oxide scale. A typical 

EDS scan in (fig. 26(c)) shows significant chemistry differences compared to the 316 °C (600 °F) C-ring specimen 

described above. S/N 120 fracture surfaces typically exhibit 

high oxygen peaks, as well as a noticeable fluorine peak, in 

addition to peaks for niobium and hafnium. 

The fractography conducted on the opened crack surfaces 

from the C-ring test specimen and S/N 120 demonstrates the 

significant differences between the laboratory-tested material 

and the flight hardware. These are important distinctions 

because the cracking in the C-ring test was used to infer crack 

initiation by hot-salt cracking in S/N 120 (ref. 33). However, 

the absence of fluorine-containing reaction products on the 

C-ring cracks does not support the hot-salt cracking mechanism. Instead, the C-ring testing provides support for hydrogen embrittlement, because the crack surfaces do not contain 

reaction products, similar to those from hydrogen-assisted 

cracking (ref. 34). Although the initial intent of these tests was 

to introduce a salt (NbF5), while eliminating all sources of 

hydrogen, the elevated temperature exposures were performed 

in laboratory air. As a result, the NbF5, which is a highly 

hygroscopic compound (ref. 32), appears to have absorbed 

water, thereby providing a source of hydrogen to affect the 

mechanical behavior of the C−103 material. Thus, the C-ring 

testing and the detailed microprobe analyses did not substantiate the root cause identified in review articles (refs. 7 and 10) 

of the RCS thruster case history. Further evidence that supports hydrogen embrittlement being the root cause of the IGC 

in the RCS thrusters will be provided in sections 2.4 and 2.5. 

Section 2.4 will present electrochemical characterization data 

showing that hydrogen production in C−103 is thermodynamically assured in the presence of HF-containing etchants. Section 2.5 will demonstrate through mechanical testing and 

resultant fractography that intergranular failure is observed in 

niobium materials when charged with appropriate levels of 

hydrogen. 

2.3.3 Summary and Conclusions From C-Ring Testing 

The C-ring test results did not support the hot-salt cracking 

mechanism, which was identified in review articles (refs. 7 

and 10) as being the root cause of the IGC in the RCS thrusters. Instead, the C-ring test specimens had fracture surfaces 

with characteristics similar to those from hydrogen-assisted 

cracking. 

2.4 Electrochemical Characterization 

Based on the results from the literature review, the IGC 

observed in the RCS thrusters was most likely due to environmentally assisted cracking that was initiated during the 

processing of these components. To better understand the 

chemical interactions that may have affected the mechanical 

behavior of these components, a series of electrochemical characterization tests were developed. These tests consisted of open 

circuit measurements and the development of polarization 
NASA/TP—2005-214053 8 

curves, which can be used to determine the electrochemical 

reactions that were likely to occur during the processing of the 

RCS thrusters when the C−103 is exposed to a specific environment. In addition, electrochemical scratch tests were used 

to interrogate the electrochemical processes that occur on a 

material as the native oxide is ruptured, which is a necessary 

process to produce environmentally assisted cracking in materials, such as niobium, that develop a native oxide. Once the 

likely electrochemical processes and the test conditions that 

create these processes are identified, exposure conditions for 

representative mechanical tests can be created. 

The critical environment appears to be produced by the 

niobium etchant used during processing. This etchant consisted of HF, HNO3, and H2O in equal proportions. Niobium 

and niobium alloys are known to be attacked by HF (refs. 18 

and 35 to 37) but are resistant to HNO3 (refs. 18 and 35) due to 

the formation of an adherent oxide film. Consequently, a 

dilute HF solution was used in the electrochemical characterization tests, instead of the niobium etchant, in order to isolate 

the effects of the HF and to simulate the dilute acid condition 

that would be anticipated to collect in occluded regions of the 

injector after rinsing with water. In these characterization 

tests, HF was diluted to a pH ≈ 3.5. A neutral solution (1 wt% 

sodium chloride (NaCl) in H2O) was also evaluated to compare the response of the C−103 material in a relatively benign 

environment to that observed in the dilute HF solution. 

All C−103 specimens were extracted from a partially 

machined injector that was manufactured during the same 

approximate timeframe as injectors that have seen service on 

the shuttle orbiters. The Ti-6Al-4V specimens were extracted 

from a flight hardware flange that was bolted to thruster S/N 

120. The Ti-6Al-4V was included in the testing to determine if 

the couple created from the Ti-6Al-4V flange contacting the 

C−103 thruster was likely to affect the electrochemical processes occurring on the exposed C−103 surfaces. 

2.4.1 Open Circuit Measurements 

Electrochemical tests were initially performed to obtain 

open circuit measurements of C−103 and Ti-6Al-4V; values of 

the open circuit potential (EOC) dictate what electrochemical 

processes can occur on each material containing a native 

oxide, excluding any outside influences. The Ti-6Al-4V component was black in color due to the production of a thick 

oxide film that developed during processing. Open circuit 

potentials were obtained for the Ti-6Al-4V with this thick 

oxide film and for Ti-6Al-4V with an air-grown film produced 

after polishing the specimen. Electrochemical testing indicated 

little difference in the open circuit potentials for the titanium 

alloy with either oxide film, and thus all testing for Ti-6Al-4V 

was performed subsequently on specimens that were polished 

to a 600-grit finish. After removal of the oxide film, specimens were immediately immersed in the selected electrochemical solution, and testing was begun. Additional tests 

were performed using a deaerated solution to determine if 

there was an effect observed from the removal of dissolved 

oxygen in the electrolyte; deaeration was obtained by bubbling 

argon through the solution for at least 30 min prior to the 

introduction of the specimen. However, it was found that the 

state of aeration had little effect on the open circuit potential 

for either material in the solutions examined. Thus, all results 

presented will be for aerated solutions. In addition to the 

measurement of open circuit for each material, couple potentials were also measured by electrically coupling samples of 

each material and immersing them into the selected solution. 

For the couple measurements, the C−103 and Ti-6Al-4V 

specimens were sectioned to expose nearly equal surface areas 

of each material to the solution. All open circuit measurements 

were made by measuring the specimen potential relative to a 

saturated calomel electrode (SCE). 

The EOC data are provided in table II for C−103, Ti-6Al-4V, 

and the C−103/Ti-6Al-4V couple in aerated solutions of dilute 

HF and of 1 wt% NaCl in H2O (neutral solution). The reversible potentials (Er) for water reduction for each of the aerated 

electrolytes are also given in table II. The values of Er represent the potential at which the reduction of water will produce 

a hydrogen fugacity of 1 atmosphere (ref. 38). For potentials 

less than Er, the fugacity will be greater than 1 atmosphere.

Comparison of the data in the table shows that the open circuit 

potentials for C−103, Ti-6Al-4V, and the couple in the dilute 

HF solution are more negative, or more cathodic, with respect 

to the Er for water reduction in the dilute HF solution. However, the open circuit potentials in the neutral NaCl solution 

are more positive, or more anodic, relative to the Er for water 

reduction in the neutral solution. This suggests that hydrogen 

is more readily produced, via the reduction of water, on the 

surface of each material in the dilute HF solution than that 

which will be produced in a neutral solution. Table II also 

shows that the EOC for Ti-6Al-4V is 0.598 V cathodic to the 

EOC measured for the C−103 in the dilute HF solution, which 

suggests that a greater amount of hydrogen can actually be 

produced on the surface of the Ti-6Al-4V than on the C−103 

surface in the dilute HF solution. When the Ti-6Al-4V is coupled to the C−103 material, the EOC is observed to lie between 

the EOC values for each separate material. This suggests that if 

the Ti-6Al-4V is coupled to the C−103 component, a greater 

amount of hydrogen production may occur on the surface of 

the C−103 material than if the Ti-6Al-4V flange were not 

present. 

2.4.2 Electrochemical Scratch Testing 

Niobium and titanium alloys produce a native oxide in the 

presence of oxygen and/or water. Consequently, electrochemical tests generally interrogate these oxide films and do not 

result in a complete understanding of the reactions that occur 

on the base material. This is of particular concern when 

addressing environmentally assisted cracking events, since the 
NASA/TP—2005-214053 9 

native oxide may be breached in service due to mechanical 

rupture or abrasion of the film and the critical electrochemical 

processes that may affect material behavior may be exacerbated when the film is ruptured. To interrogate such a situation, C−103 specimens were prepared in metallographic 

mounts. One surface of the mounted specimen was polished 

through 600 grit to provide a fresh surface of the C−103 coupon. Conductive wires were spot welded to the back of C−103 

specimens. The polished surface was then coated with a nonconductive film and immersed in either the dilute HF or the 

1 wt% NaCl solution. The location of the spot weld was such 

that it would not be exposed to the electrolyte. The nonconductive film and native oxide were then ruptured with a diamond scribe to expose fresh base metal to the solution. A 

high-speed data logger (250 Hz) was used to acquire potential 

versus time data during the scribing of the surface. By measuring the potential transient created, a better understanding can 

be obtained of the electrochemical processes occurring at the 

crack tip as the oxide film is ruptured during crack growth. 

The potential transients obtained from the scratch tests of 

the C−103 material in the dilute HF solution are shown in 

figure 27. Because this test measures a rapid transient in 

behavior, a minimum of three measurements was performed 

for each test condition to ensure that an accurate representation of the electrochemical process was obtained. The average 

scratch potentials for C−103 immersed in the dilute HF and 

the neutral NaCl solutions are given in table III; scratch potentials (Escratch) represent the minima observed in the curves in 

figure 27. The Escratch for C−103 were found to be more negative, or more cathodic, compared to the EOC for the C−103 in 

the same electrolyte. The measured scratch potentials suggest 

that hydrogen can readily form on the fresh surface of C−103 

as the oxide film is ruptured. Hydrogen production can be 

quantified by calculating the hydrogen fugacity (fH2

) for a 

given potential, as shown in equation (1) (ref. 38). 

 2log H 30

E E r f ⎛ ⎞ − ⎜ ⎟ = − ⎝ ⎠ (1) 

where E is either the measured EOC or the measured Escratch, 

and Er is the reversible potential for the reduction of water. 

Values of potential in equation 1 are in units of millivolts 

(mV) and fugacity is in atmospheres. The fugacities calculated 

for the scratch potential and open circuit potentials for C−103 

in the dilute HF and 1 wt% NaCl are also provided in table III. 

The resulting fugacity at a bared C−103 surface is shown to be 

more than 15 orders of magnitude greater that on the native 

oxide surface for either solution. 

In summary, it has been shown that the open circuit potentials for both the Ti-6Al-4V material and the C−103/Ti-6Al4V couple are cathodic in the dilute HF solution, with respect 

to the open circuit potential for the C−103 material. The couple potential suggests that cathodic processes and hydrogen 

production can occur on the C−103 surface in the presence of 

the Ti-6Al-4V flange that may not occur without this flange or 

at distances sufficiently far away from this flange to be unaffected by the couple. However, when the native oxide is ruptured, the electrochemical processes in both the dilute HF and 

the NaCl solutions are highly cathodic, relative to the processes operative with an intact native oxide. Furthermore, the 

measured scratch potential is similar in sign and magnitude to 

that of the open circuit potential for the C−103/Ti-6Al-4V 

couple. Consequently, if the native oxide becomes ruptured in 

service, the Ti-6Al-4V flange does not need to be present nor 

involved in the electrochemical processes for hydrogen production to occur and affect the environmentally assisted cracking in the C−103 thrusters. 

2.4.3 Electrochemical Polarization 

To determine the overall electrochemical behavior for the 

C−103 material, polarization curves were generated. By 

evaluating the electrochemical response over a large range in 

potential, a better understanding of how electrochemical processes are affected by outside influences, such as coupling a 

piece of C−103 to Ti 6Al-4V or rupturing the oxide film from 

a niobium alloy, can be determined. Here, two measurements 

(one cathodic and one anodic with respect to the open circuit 

potential) were acquired for each solution in the aerated condition. The tests consisted of polishing a specimen to a 600-grit 

finish, exposing the specimen to the selected solution, allowing an open circuit condition for 1 hr, and then performing a 

controlled potential scan at a rate of 0.167 mV/sec, while 

monitoring the current. For each anodic scan, the initial potential was set to be 40 mV cathodic to the open circuit potential, 

while the cathodic scan was initiated at 40 mV anodic to the 

open circuit potential. These starting conditions for the electrochemical scans made it possible to have a small section of 

data from each test under the same conditions, which was used 

to validate the results. 

The polarization curves acquired in the dilute HF and the 

neutral (1 wt% NaCl in H2O) solutions are compared in figure 28. Here, the open circuit potentials (approx. −0.4 to 

−0.6 VSCE in fig. 28) are approximated by the lowest measured 

current values. All data at higher values of potential indicate 

anodic processes (i.e., dissolution and oxidation), while all 

data at lower values of potential indicate cathodic processes 

(i.e., water reduction). A large passive region is shown to start 

with a “knee” in the data at approximately −0.25 VSCE, and the 

passive region remains stable with no significant increase in 

current for more than 2 V in potential, suggesting a protective 

passive film is present over a large range of potential. However, it should be noted that the current density throughout the 

passive region is more than 1 order of magnitude higher in the 

HF solution than for the neutral solution, indicating significant 

anodic processes are occurring throughout the passive region. 

Comparison of the two polarization curves in figure 28 also 
NASA/TP—2005-214053 10 

reveals a higher exchange current density (io) and a lower EOC

are observed in the HF solution compared to that in the neutral 

solution. This indicates that the C−103 material is more electrochemically active in an HF solution than in the neutral 

aqueous solution. In the cathodic region in figure 28, the current density is slightly higher in the HF solution; however, the 

shapes of the two curves are similar, suggesting similar 

cathodic processes (i.e., water reduction) will occur in either 

of the two solutions. 

The polarization curve generated for the C−103 exposed to 

the dilute HF solution has been isolated in figure 29 for further 

discussions. The current values cathodic to the EOC are 

observed to be much higher than in the anodic region, suggesting a larger rate of reaction for cathodic reactions compared to 

the anodic reactions. The values of the open circuit and scratch 

potentials have also been superimposed in figure 29. These 

results indicate that at the open circuit (at least near the titanium flange) and scratch potentials in this system, cathodic 

processes are expected to dominate the electrochemical processes of this system. This observation implicates the role of 

hydrogen in the environmentally assisted cracking processes 

observed for the thruster components. These electrochemical 

tests clearly show that atomic hydrogen is readily produced on 

the bared C−103 surfaces exposed to aqueous solutions. Additionally, the relatively high cathodic currents in figure 29 

indicate that the amount of hydrogen that can be produced via 

water reduction may be significant. While these characterization tests demonstrate that atomic hydrogen is readily formed 

on the surface of the C−103 material, the tests do not shed any 

light on how this hydrogen production may affect the C−103 

material. For example, it would be instructive to determine if 

the hydrogen being produced readily adsorbs onto the C−103 

surfaces, then absorbs into the material and diffuses into the 

bulk (ref. 39). Therefore, mechanical screening tests were 

subsequently performed to determine how the C−103 material 

behaves when it is charged with hydrogen. 

2.4.4 Summary and Conclusions for Electrochemical 

Characterization 

Hydrogen can readily form on the surface of the C−103 

thrusters when exposed to an HF-containing etchant. The 

hydrogen is produced via the water reduction reaction. The 

driving force for hydrogen production is very high on the 

surface bared of the native oxide film. 

(1) The presence of HF increases the electrochemical activity of C−103. 

(2) C−103 is not susceptible to rapid dissolution, even in 

dilute HF solutions. However, the presence of dilute HF can 

result in increased electrochemical processes. 

(3) The production of H is thermodynamically assured on 

the surface of bared C−103 when moisture is present. 

(4) Coupling Ti-6Al-4V to C−103 may increase the production of hydrogen in dilute HF relative to the amount of 

hydrogen produced on C−103 without titanium being present. 

(5) The hydrogen fugacity on a bared C−103 surface in 

dilute HF is ~2×1015 atm. 

2.5 Environmentally Assisted Cracking (EAC) 

Screening Tests 

A series of uniaxial tension and bend tests were designed to 

determine how the exposure to environment and temperature 

can affect the mechanical behavior of the C−103 material. All 

uniaxial tension specimens were run by exposing the specimen 

to a dilute etchant solution and applying a constant potential. 

Each potential was selected based on the electrochemical tests 

presented in section 2.4. By applying a constant potential to 

the test specimen, the electrochemical processes for each test 

can be controlled so that specific damage processes can be 

evaluated. These tests are designed to screen specific environmentally assisted cracking processes (i.e., dissolutionbased SCC or hydrogen embrittlement). Once these processes 

were understood, bend specimens were designed that more 

accurately reflect the loading and exposure conditions observed in service. Mechanical test specimens were obtained 

from a partially machined injector (see fig. 30) that was manufactured during the same approximate time frame as injectors 

that have seen service on the shuttle orbiters. 

2.5.1 Uniaxial Tension Testing 

Uniaxial tension specimens were machined with a gage section measuring approximately 5.1 by 5.1 mm (0.2 by 0.2 in.), 

as seen in figure 31. The specimens were notched with a brass 

wire via electron discharge machining (EDM) in order to 

provide a local stress concentration at the midsection of each 

specimen. The notch was machined across the width of one 

face and measured between 0.5 and 0.63 mm (0.02 and 

0.025 in.) in depth. Each test specimen was placed in an electrochemical cell with the gage section submersed in a dilute 

HF solution (diluted to a pH ≈ 3.5), while the threaded specimen ends were outside of the solution and placed into an adaptor of a load frame. With this testing setup, the specimen gage 

section could be within an electrochemical cell while being 

stressed under constant load. A reference and platinized 

counter electrode were also immersed within the solution of 

the electrochemical cell. The reference electrode, counter 

electrode, and specimen were connected to a potentiostat to 

allow each test to be performed under a constant potential 

condition. Durations of exposure were selected at 24, 48, and 

~90 hr, although some specimens failed during exposure under 

applied load. A list of testing conditions is provided for each 

specimen in table IV. Due to the low conductivity of the test 

solution, 1 wt% NaCl was added to the test solution initially to 
NASA/TP—2005-214053 11 

ensure that the desired electrochemical potential was in fact 

the potential generated at the surface of the test specimens. 

There were no significant differences between the tests conducted with or without the NaCl. 

The range of potentials examined was based on the electrochemical tests described in the previous section (see fig. 29). 

As seen in table IV, several tests were performed with the 

applied potential between −600 mVSCE and −1 VSCE. These 

values of potential were selected to interrogate the effect of 

cathodic processes on the mechanical behavior of C−103. The 

scratch potential should be the most cathodic potential established on the surface of the thruster material, unless it is affected by an outside influence not considered in this study. 

Thus, initially tests were performed at −900 mVSCE, which is 

slightly cathodic to the measured scratch potential in the dilute 

HF solution (see fig. 27). Then the potentials were increased 

toward EOC until little to no embrittlement was observed. At 

−600 mVSCE, there was limited embrittlement, and as a result, 

no additional cathodic potentials were selected near this range. 

However, an additional test was performed at −100 mVSCE to 

examine the potential for aqueous dissolution-based SCC. The 

potential of −100 mVSCE was determined based on the anodic 

polarization of C−103 (see figs. 28 and 29). Here, −100 mVSCE

is slightly anodic to the passive knee, which is considered to 

be a region that is highly susceptible to SCC (ref. 40). In addition to examining the influence of the etchant solution, the 

potential effects of the thermal bakeout used in the manufacture of the thruster components (316 °C (600 °F) for 48 hr) 

was studied by removing several of the exposed specimens 

from the electrochemical exposure cell and placing them into 

air furnace under load. Upon completion of these tests, all 

specimens were strained to failure, and examination of the 

fracture surfaces was subsequently conducted using SEM. 

Two of the test specimens (specimens 1 and 2) failed during 

the exposure process, as indicated in table IV. These test 

specimens exhibited highly embrittled fracture surfaces with 

little observed ductility, an example of which is given in figure 32. Smooth terrace features along with visible slip bands 

and fluted features are clearly visible in figure 32(b), indicating a predominantly cleavage fracture. Although there was 

some IGC visible on this fracture surface, the predominant 

fracture process was transgranular cleavage, which is distinctly different from the intergranular fracture observed in 

service. Specimen 2 was statically loaded to approximately 

177 MPa (25.6 ksi), with an applied potential of −750 mVSCE, 

and failed after only 26.5 hr. Similar failures were observed 

for specimen 1, which had an applied stress of approximately 

198 MPa (28.8 ksi) and an applied potential of −900 mVSCE; 

failure occurred in this specimen in less than 40.5 hr. As the 

exposure time was decreased for the same applied potential 

(specimen 3 in table IV) or the potential was increased to more 

anodic potentials (−600 mVSCE for specimen 4 in table IV), the 

fracture surfaces suggested that the outer specimen surfaces 

were less ductile than the central portion of the specimens. 

However, there were no indications of intergranular or cleavage fracture, thereby indicating only a moderate embrittlement 

was induced from the exposure process for specimens 3 and 4.

It should be noted that specimen 4 was also baked out prior to 

mechanically overloading the specimen. This bakeout was not 

performed for specimens 1 through 3, and the potential effects 

of the bakeout process will be discussed below for specimens 

5 through 9. Although the distinct change in fracture mode 

observed for specimens 1 and 2 suggest that these specimens 

were readily embrittled by the absorption of hydrogen into the 

C−103 material, the cleavage fracture morphology produced 

in these tests did not resemble the intergranular fracture produced in service. Consequently, additional tests were performed to fully characterize the damage processes experienced 

during the processing of the cracked thruster components.

Five test specimens were examined by exposing the specimens to the dilute HF solution at −1 VSCE for 90 hr either 

using a small tensile load (specimens 5 and 6) or zero load 

during exposure (specimens 7, 8, and 9), as seen in table IV. 

Predominantly brittle fracture was observed when these 

specimens were strained to failure following exposure (specimens 5, 7, and 8), while predominantly ductile fracture was 

observed (specimen 6) if the exposure was followed by a heat 

treatment of 316 °C (600 °F) for 17 hr. Specimen 9 had a 

relatively small region (~0.25 mm (0.01 in.)) around the outer 

periphery of the specimen gage that exhibited 100 percent 

IGC, while the majority of the fracture surface was ductile 

(see fig. 33). These features on the fracture surfaces suggest 

that hydrogen was absorbed within the bulk of the specimens 

during the −1 VSCE exposures and that the hydrogen was 

highly mobile in the C−103 material even at room temperature and with no external loading. Additionally, the intergranular features on specimen 9 suggest that at 316 °C 

(600 °F) the absorbed hydrogen can readily diffuse and repartition within the C−103 microstructure, resulting in intergranular fracture. 

To interrogate the potential role of SCC in the damage of 

the C−103 material, a specimen (specimen 10 in table IV) was 

loaded to 221 MPa (32 ksi) and exposed to the dilute HF solution for more than 90 hr at an applied potential of 

−100 mVSCE. The potential of −100 mVSCE was chosen to 

represent the most susceptible condition for SCC processes in 

the dilute HF solution, based on the electrochemical testing 

presented previously in section 2.4.3. Following the 90 plus 

hours of exposure, uniform corrosion of the surface was 

observed on the surface of the specimen (see fig. 34); however, no cracking was observed. 

These results show that the RCS thruster material can suffer 

from hydrogen embrittlement when exposed to a hydrofluoric 

acid containing etchant similar to that used during the processing of the RCS thruster. It should be noted that the electrochemical characterization tests presented earlier in section 2.4 
NASA/TP—2005-214053 12 

also indicate that hydrogen can be produced under certain 

conditions with C−103 in water, although the solution is not as 

electrochemical active without the HF present. However, the 

uniaxial tension testing in the dilute HF solution did not 

reproduce the considerable IGC observed in the thruster material. This suggests that while the absorption of hydrogen 

within the bulk of the C−103 is affecting the cracking processes in the thrusters in service and in these uniaxial test 

specimens, the absorbed hydrogen level and/or loading conditions used in these tests were not sufficient to replicate the 

exact fracture mode observed in service. A change in fracture 

mode for hydrogen embrittlement processes based on 

absorbed hydrogen level and loading condition has been previously reported for other material systems (refs. 19 and 41). 

2.5.2 Bend Testing 

A total of four bend specimens was examined to more accurately reflect the exposure of etchant solution and the 

mechanical loading (i.e., bending versus uniaxial tension) 

observed in the relief radius region for the thruster components. The procedures used in this testing were based on previous tests performed at Johnson Space Center (ref. 42) in 

collaboration with this study. Rectangular test specimens 

having a minimum thickness of 2.5 mm (0.1 in.) were 

machined. In the middle of one broad surface of the specimen, 

an indentation was made using a Rockwell ball indenter. As 

illustrated by the schematic in figure 35, each specimen was 

then placed in a three-point bend fixture such that a tensile 

stress of 276 MPa (40 ksi) was applied at the indentation; the 

stress was calculated without consideration of the reduced 

cross-sectional area at the indentation and without consideration of any change in the residual stress produced from the 

indent. The applied stress was selected to be consistent with 

the stress calculated (ref. 33) for the relief radius in previous 

studies and approached the reported values for the 0.2-percent 

yield strength (ref. 35) for C−103. The niobium etchant (HF, 

HNO3, and H2O in equal proportions) was then diluted by 

50 percent, and a few drops of it were placed on the specimen 

to cover the indentation mark. A sheet of Ti-6Al-4V was then 

placed over the indentation, and the entire specimen was 

sealed in a cell to maintain the air at a relative humidity (RH) 

in excess of 70 percent, thereby limiting evaporation of the 

diluted etchant. Each specimen was exposed to this condition 

for a minimum of 36 hr. The environmental cell was then 

removed, and each specimen was baked in air for 48 hr at 

316 °C (600 °F), while maintaining the applied bending load. 

After each specimen cooled to room temperature, wires were 

spot welded in close proximity to the indentation on each 

specimen, in order to monitor crack growth using a direct 

current potential difference (DCPD) method (ref. 43). After 

wire attachment, the environmental cell was reintroduced 

to expose each specimen again to a moist air environment 

(>70-percent RH), with exposure times varying from 7 to 

45 days. Following this exposure process, each specimen was 

overloaded in bending to reveal any cracking that may have 

occurred during the environmental exposure. 

Significant cracking (>0.36 mm (0.014 in.) in depth) was 

observed in each of the four bend specimens, and the resultant 

fracture surfaces were 100 percent intergranular in nature, as 

shown in figure 36. In addition to the intergranular nature of 

the cracking, “popcorn” features were evident on the fracture 

surfaces (figs. 37(a) and (b)). These popcorn features are consistent with those observed on the intergranular cracks in the 

RCS thrusters, and details of these observations in S/N 120 

and 132 will be described in sections 3.1 and 3.2, respectively. 

The similarities in the microstructural observations between 

the test specimens and in-service hardware suggest that the 

simulated damage processes are consistent with those produced during the processing of the thruster components that 

have experienced environmental cracking. 

The DCPD measurements did not resolve any crack extension during the high humidity exposures following bakeout, 

even for the specimen tested for 45 days. Figure 37 contains a 

series of micrographs at the tip of one of the intergranular 

cracks produced in bend specimens. The popcorn features 

observed on the fracture surface of the test specimens are 

found in abundance across most of the environmental crack 

(figs. 37 (a) and (b)). Only a small region (i.e., five grains in 

length at the crack tip) of relatively clean grains have been 

observed in each specimen studied (see fig. 37(c)). This 

observation is consistent from specimen to specimen, is independent of exposure time, and is consistent with observations 

for cracks produced in S/N 132. Most importantly, these 

observations suggest that the IGC produced in each specimen 

was present prior to the high humidity exposures, and significant crack growth (greater than 0.08 mm (0.003 in.)) will not 

occur in service following the bakeout process without the 

introduction of additional etchant. 

2.5.3 Summary and Conclusions for Mechanical Testing 

The IGC observed in the RCS thruster components is due to 

hydrogen embrittlement. When exposed to a dilute HF solution, a bending load, and elevated temperature (316 °C 

(600 °F)), IGC consistent with that observed in service has 

been produced. 

(1) Dissolved hydrogen can result in various fracture 

modes for C−103, which are dependent upon the absorbed 

hydrogen level and loading condition. 

(2) C−103 does not suffer from dissolution-based SCC 

under the range of conditions examined in this study. 

(3) No significant crack propagation is expected to occur in 

service for the RCS thrusters, even under exposures to moist 

air environments after the elevated temperature bakeout, as 

long as additional HF-containing compounds are not introduced to the components. 
NASA/TP—2005-214053 13 

Section 3: Fractography of RCS Thruster 

Hardware 

The first failure analysis conducted on RCS thruster S/N 

120 was performed at Boeing HB in June 2004, the results of 

which were described previously in section 1. The fracture 

surfaces that were opened and examined at Boeing HB, along 

with the remaining pieces after machining the injector half 

shown in figure 4, were supplied to GRC for subsequent 

examination by the NESC Materials SPRT. The purpose of 

this additional fractography by the NESC Materials SPRT was 

to verify fracture damage modes on the actual hardware and to 

employ additional microscopy techniques in an effort to look 

for evidence of crack propagation. 

3.1 Observations in S/N 120 

The failure analysis on S/N 120 was initially conducted on 

the half of the injector (fig. 4) that included bolt holes H 

through O. The other half of S/N 120 that included bolt holes 

P through G was set aside for NDE development, although it 

was used later in May 2005 to investigate repair and rejuvenation processes at WSTF. A schematic diagram of S/N 120 in 

figure 38 identifies the approximate thruster locations from 

which crack depth data were obtained, as well as the organizations that produced these data. Ten out of sixteen counter 

bores (A through P) were cracked in S/N 120, and these are 

indicated in the figure by the filled red circles; unfilled circles 

indicate that the counter bores were uncracked. For the relief 

radius cracking, dye penetrant inspection indications from 

Boeing HB (ref. 4) and WSTF are shown in the schematic by 

the blue and green lines, respectively. The former was 

obtained in 2004 on the entire thruster circumference, whereas 

the latter was obtained in 2005 only on the injector half that 

included bolt holes P through G. Dye penetrant inspections 

obtained at WSTF indicated more extensive cracking than 

those from reference 4. 

The crack depth data are summarized in figure 39 as a function of angle around the circumference of the thruster; the blue 

diamonds in the figure representing the relief radius cracking 

and the pink squares representing the counter bore cracking. 

More data have been obtained for the depth of relief radius 

cracking than that for the counter bore cracking. Figure 39 

indicates significant cracking around the circumference, 

although the depths of cracking are highly variable. No apparent relationship was evident between the depth of the cracking 

and the position around the circumference. Assembly stresses 

are believed to contribute to the relief radius cracking as a 

result of the bolting together of nonflat flanges. The cracking 

in S/N 120 is relatively shallow compared to another thruster, 

S/N 132; crack depth profiles will be presented for S/N 132 in 

section 3.2. It is unknown how these crack depths compare to 

those in the thruster fleet. 

Relief radius cracks in S/N 120 were opened up to view the 

intergranular fracture surfaces in order to confirm that the 

cracking was produced during manufacturing in the presence 

of fluoride-containing acids. Scanning electron microscopy, 

EDS, and Electron Spectroscopy for Chemical Analysis 

(ESCA) were performed on the opened crack faces. The reaction product chemistries, morphologies, and densities were 

examined from the crack mouth to the crack tip for the following portions of S/N 120: the relief radius cracking between 

bolt holes I and J; the relief radius cracking between bolt holes 

J and K; and the relief radius cracking near bolt hole L. Similar results were obtained from each of these areas; much of the 

data that will be presented was obtained from the relief radius 

cracking between bolt holes I and J, unless otherwise noted. 

These bolt holes are identified in figures 4 and 38. 

3.1.1 Relief Radius Cracking Observations. 

The cracking on the relief radius surface was observed in 

the SEM prior to the lab opening of the cracks. Figure 40(a) 

shows the relief radius surface and its proximity to bolt hole J. 

It is evident from figure 40(b) that the relief radius cracking is 

not a single crack but a series of multiple cracks with separate 

initiations. Many cracks appear interconnected and branching 

is evident, whereas other cracks are observed to overlap without joining. The relief radius cracking between bolt holes I and 

J was subsequently opened and revealed an intergranular fracture surface with no change in fracture mode observed at the 

crack tip. An overall image of the intergranular surface is 

shown in figure 41(a); the deepest part of the intergranular 

crack in this area was measured to be 2.5 mm (0.097 in.). A 

deeper crack (5.0 mm (0.195 in.)) was also opened 19° away 

in a clockwise direction around the circumference near bolt 

hole J and is shown in figure 41(b); this montage of BSE 

images reveals that the cracking occurs over slightly different 

elevations as a result of the separate crack initiations at the 

relief radius. The intergranular surface is oxidized across its 

depth, as indicated by the darker appearance of the intergranular crack relative to the lighter lab fracture, which is freshly 

exposed C−103. Although the major crack direction is oriented vertically in the montage in figure 41(b), locally the 

cracking may progress in different direction as cracked ligaments link together; this observation is supported by the 

appearance of lighter areas (circled in fig. 41(b)) within the 

major intergranular cracked region. 

The higher magnification in figure 41(c) reveals more 

clearly the intergranular nature of the fracture surface, as well 

as the presence of products on the grain facets. Although predominantly intergranular, the fracture surfaces contain secondary cracks that are transgranular (fig. 41(d)). These 

transgranular cracks were also present in polished mounts of 

other sections of the relief radius cracking in S/N 120, which 

demonstrates that the secondary cracks were not associated 

with lab opening procedures. Most of the transgranular cracks 
NASA/TP—2005-214053 14 

observed were very shallow, as seen in the polished section in 

figure 41(e). Microprobe line scans were performed traversing 

across the intergranular crack in figure 41(e), into the adjacent 

regions containing the secondary cracking, and into the grain 

interior. No chemical differences were detected between the 

regions containing the transgranular cracking and the neighboring C−103 material. Thus, the transgranular cracks extended 

beyond the niobium oxide and reaction products on the intergranular surfaces and penetrated into the C−103 parent material. 

Secondary cracking may be a signature of hydrogen-assisted 

cracking. 

It is evident from the representative views in figures 41(c) 

and (d) that “bumps” and cracked nodules were prevalent on 

the intergranular crack surfaces. The cracked nodules resembled opened popcorn kernels that appeared to be residing on 

top of the intergranular fracture surface. Some grain facets had 

higher densities of the popcorn than others, but nonetheless 

popcorn is observed from the crack mouth to the crack tip. 

Quantitative analyses of the popcorn density showed no distinct changes in the popcorn density on S/N 120 crack surfaces 

as a function of distance to the crack tip; these data will be 

discussed in detail later in section 3.3 in this report. EDS scans 

were obtained from various features on the intergranular surface at the mouth, tip, and the midpoint of the crack. Figure 42(a) shows a secondary electron image of the crack tip 

and the adjacent overload area ahead of the crack, and figure 42(b) is the corresponding BSE image. Secondary electron 

images provide more topographical information about a given 

area, whereas BSE images additionally provide indications of 

atomic number differences in the features within the field of 

view. EDS scans were obtained from locations labeled “C,” 

“D,” and “E” in figures 42(a) and (b); most of the EDS scans 

were taken at an accelerating voltage of 6 keV, which enables 

more of a surface analysis to be performed because the characteristic x-rays generated are confined to near-surface depths. A 

representative EDS scan of the popcorn nodule, labeled “C,” is 

given in figure 42(c); the popcorn was found to contain high 

oxygen and fluorine peaks, in this case with roughly a 2:1 

peak ratio, respectively. A peak for carbon was also exhibited 

adjacent to the niobium M line; and niobium (L line) and 

hafnium peaks are also evident in the EDS scan. Numerous 

popcorn nodules were analyzed across the fracture surface, 

and very similar results were obtained. The popcorn was identified as most likely being NbO2F, based on semiquantitative 

analyses, which is an expected reaction product that forms 

after NbF5 is exposed to moisture in the air (ref. 32) at elevated temperatures. It is not clear if the carbon evident in the 

EDS scans is tied up chemically in the popcorn or if carbon 

was introduced as a contaminant after popcorn formation. 

Single crystals of niobium have been shown in the literature 

to form oxide nodules (ref. 44) that have a very similar 

appearance to the popcorn observed on the S/N 120 crack 

surfaces. Pure niobium initially forms a protective oxide film 

that has a rippled or bumpy texture; the film becomes nonprotective as oxidation proceeds due to a breakaway oxidation 

phenomenon, which is characterized by the formation and 

growth of isolated patches of locally thick oxide. Small cracks 

form in the locally thick oxide scales, which enable more 

oxygen to diffuse into the niobium substrate beneath the scale. 

This significantly increases the rate of oxidation beneath these 

cracked areas and promotes the formation of nodules or popcorn at these locations. Reference 44 also demonstrates that 

crystallographic orientation influences the rate of oxidation in 

single crystal niobium, which is consistent with the observation in S/N 120 that some grain facets have more popcorn 

formation than others. It will be shown later in section 3.3 that 

popcorn does form during the oxidation of C−103 at thruster 

manufacturing bakeout conditions (48 hr in air at 316 °C 

(600 °F)) but only if fluorine is present. 

The oxide scale on the intergranular surfaces near the popcorn was also analyzed for comparison to the popcorn. The 

EDS scan in figure 42(d) for location D (in fig. 42(a)) shows 

high oxygen and similar niobium and hafnium peaks compared to the popcorn. Although fluorine was detected in the 

oxide scale, its peak height was typically much lower than 

those for oxygen and niobium (L line) in the oxide scale. The 

fluorine content in the oxide scale is also lower than that in the 

popcorn. It should be noted that the carbon peak was not evident at location D, and this was a typical result for the light 

regions on the oxidized intergranular surface in figure 42(b). 

However, the EDS scans were significantly different when 

obtained from the darker regions (location E) on the fracture 

surface of figure 42(b); darker regions in BSE images indicate 

the presence of lower atomic elements. In these areas, the 

carbon peak was one of the strongest peaks observed; sodium 

and magnesium peaks were also exhibited (fig. 42(e)). The 

presence of sodium is often considered to be a contaminant 

that may be the result of handling issues. 

Other dark regions were also observed with BSE imaging, 

especially along changes in elevation seen on the fracture 

surfaces, such as in valleys created between grain facets as 

seen in figure 43(a). Dominant carbon and fluorine peaks were 

observed in the EDS spectra for these dark areas (figs. 43(b) 

and (c)). The carbon- and fluorine-rich material did not appear 

to be tied up in the oxide because the material was observed to 

vaporize under the electron beam in the SEM. Subsequent 

analyses conducted on the cutting fluid used in sectioning this 

sample prior to the lab opening show a distinctly different 

chemistry from that shown in figures 43(b) and (c). It is possible that the dark carbon- and fluorine-rich material on the 

fracture surface is residual Krytox (DuPont, Wilmington, DE) 

or residual dye penetrant material. Krytox is believed to have 

been used as a lubricant for the counter bore bolts, and detailed ESCA analyses (section 3.4) confirmed that a compound 

with a signature similar to Krytox was present on the intergranular surfaces of S/N 120. Krytox is reported to react with 
NASA/TP—2005-214053 15 

niobium oxides (ref. 45); subsequent reactions can cause the 

release of trace amounts of HF (ref. 46), which could result in 

IGC in a manner similar to that which is believed to have 

occurred with the preweld etchant during the 316 °C (600 °F) 

thruster bakeouts. 

Other morphological features were also present in isolated 

areas on the intergranular crack surfaces. Examination near the 

crack mouth (fig. 44(a)) revealed the presence of cubic-shaped 

products and some regions that looked like “mud-flats.” EDS 

scans in figures 44(b) and (c) show that both of these areas 

were also rich in oxygen, fluorine, carbon, and niobium. 

Strong sodium peaks were also evident in the mud-flat areas. 

The cubic-shaped and mud-flat products were confined to 

localized areas in the crack mouth. Acicular morphologies, 

rounded particles, and clusters were also observed across the 

intergranular fractures; examples of these features are shown 

in figure 45. The acicular products may be oxides, since EDS 

indicated large oxygen peaks, although other reaction products, or contaminants and debris that may have been swept 

into existing cracks could not be ruled out. The small surface 

bumps, uncracked rounded particles, and clusters are likely 

different variants of popcorn nodules in various stages of 

development. The EDS scans confirm the microprobe data, 

which indicate that oxygen, fluorine, and carbon are clearly 

present on the intergranular cracks of S/N 120. This is also 

consistent with scanning Auger results (refs. 4 and 8) from the 

crack surfaces and machined injector surfaces in which higher 

concentrations of fluorine, carbon, and oxygen were found for 

a 1 to 2 μm depth before returning to nominal levels. 

The opening of the intergranular crack also enabled a depth 

profile to be obtained through the oxide scale in a uniquely 

positioned grain on the fracture surface. The intergranular 

surfaces were cracked nearly all the way around this particular 

grain, except for one remaining ligament that kept the grain 

attached to the opened fracture surface (fig. 46(a)). During the 

lab opening, one of the grain facets fractured, which revealed 

fresh C−103 in an orientation that was normal to the electron 

beam in figure 46(b); the freshly exposed C−103 is lighter in 

color than the oxidized scale on the intergranular surface. The 

positioning of the grain provided a longitudinal cross section 

of the surface products and enabled a depth profile to be 

obtained through the oxide scale from the top intergranular 

surface down to the C−103 substrate. Figure 46(c) shows a 

higher magnification view of the oxide scale and C−103 substrate; D, E, and F represent the locations for the EDS scans 

shown in figures 46(d), (e), and (f), respectively. Comparison 

of the EDS data shows that the top layer in figure 46(d) contained the highest fluorine level. Lower fluorine levels are 

evident at location E, which is approximately midway between 

the oxide surface and the C−103 substrate, and the EDS scans 

(fig. 46(e)) from location E were dominated by the oxygen and 

niobium peaks. The material just under the oxide scale in 

location F appeared to be consistent with baseline C−103 that 

had been slightly oxidized (fig. 46(f)); a small oxygen peak is 

evident but the fluorine peak is not above background levels. 

Thus, it appears that the fluorine content decreases through the 

depth of the surface product as the niobium base alloy was 

approached. 

These depth profile results by EDS were repeated in three 

other locations, and the observed trends were very consistently 

repeated. In addition, an x-ray line scan was obtained from the 

surface of the oxide scale through to the freshly exposed 

C−103 in increments of 0.02 μm. Figure 47(a) shows that the 

oxygen levels are highest in the middle region, after which the 

oxygen drops off dramatically at the C−103 substrate. The 

fluorine levels are again highest near the surface and decrease 

through the thickness of the scale. Figure 47(b) shows that the 

carbon was more heavily concentrated near the surface of the 

oxide scale, which is consistent with scanning Auger results 

(refs. 4 and 8). Thus, a composition gradient is clearly present 

through these oxide scales with the fluorine content decreasing 

through the depth of the oxide as the C−103 is approached. If 

the IGC was caused by a dissolution process, instead of by 

hydrogen embrittlement, then the fluorine levels would be 

expected to be highest in the regions adjacent to the niobium 

parent material. Therefore, these data also do not support a 

dissolution process by stress corrosion or hot-salt cracking. 

3.1.2 Conclusions and Summary of Observations for 

S/N 120 

Cracking was intergranular from the mouth to the tip with 

no change in crack mode observed. No evidence was found for 

crack propagation after manufacturing. 

(1) The intergranular crack faces contained significant oxidation with popcorn oxidation evident from the mouth to the 

tip. Although some grain facets exhibited more popcorn formation than others, quantitative analyses of the popcorn density showed no distinct changes in popcorn density as a 

function of distance to the crack tip. 

(2) Fluorine, carbon, and oxygen were readily detected on 

the crack surfaces. Depth profiling through the reaction product scale indicated that the highest concentrations of fluorine 

were found at the surface and that the fluorine content diminished as a function of distance as the niobium parent material 

was approached. 

(3) Relief radius cracking was observed to be not a single 

crack but a series of interconnected cracks with separate initiations. The depth of cracking was highly variable. 

3.2 Observations in S/N 132 

3.2.1. Background 

Ultrasonic (UT) NDE inspection of RCS thruster S/N 132 

on April 12, 1979, revealed relief radius cracks (ref. 47). A 
NASA/TP—2005-214053 16 

partial history of S/N 132 is also documented in reference 48. 

After the cracks were discovered, S/N 132 was selected for 

use as a qualification test article. Ultrasonic inspections were 

performed after each round of testing with no crack growth 

detected, although it was not known at the time that this NDE 

method was unreliable. The thruster was reassembled in 

March 1982 with new insulation and was given a second 48-hr 

insulation bakeout in air at 316 °C (600 °F) (ref. 47). The 

thruster experienced an unstable operation during a test firing 

on March 23, 1983 (ref. 47), which resulted in injector temperatures exceeding 260 °C (500 °F) for a 24 sec duration with 

a 304 °C (580 °F) peak temperature reached (refs. 47 and 48). 

On August 4, 1983 (ref. 47), S/N 132 failed an external leak 

check, which indicated that the cracks had propagated through 

the injector wall to the acoustic cavity(ies). The unit was used 

in 2004 and 2005 for the development of NDE methods for the 

reliable detection and quantification of relief radius cracks. In 

addition, two small pieces of this injector were examined by 

Boeing HB in 2004 to determine if, and by what mechanism, 

the cracks propagated during thruster qualification testing; 

these results were reported in reference 48. The report concludes from the two pieces (labeled acoustic cavity (AC) 

numbers 14 and 16) examined that the intergranular crack 

propagated from the relief radius to the acoustic cavities by 

the same failure mechanism with no change in failure mode. 

The cracks were believed to have propagated during the original failure event with no evidence found for intermittent, timedelayed failure or for fatigue. The report also concludes that 

hydrogen embrittlement did not occur because of the presence 

of oxygen- and fluorine-bearing deposits on the intergranular 

fracture surface. 

The NESC strongly recommended to the Orbiter Project 

Office that a destructive examination of another thruster be 

performed by the NESC Materials SPRT. The purpose of the 

examination was twofold: (1) to determine if there was any 

evidence for fatigue crack growth during thruster operation 

and (2) to define the crack front for input to the NESC Stress 

and Fracture SPRT in their fatigue and fracture modeling. The 

crack depth profile of the relief radius cracks and the symmetry of cracking would be defined quantitatively, in order to 

benchmark the starting crack profile for vibration load analyses. RCS thruster S/N 132 was supplied to GRC in March 

2005 as a result of this recommendation. The two large sections (a 57° and a 203° piece) of S/N 132 that were destructively analyzed are shown in figure 48. The extent of this 

examination enabled the most comprehensive data to be 

obtained to date on crack depth profiles for RCS thruster 

hardware. 

3.2.2 Crack Surface Observations 

Replicas were obtained for future examination from the 

entire relief radius and all counter bores (half of N to P and C 

to K) from the two large sections prior to further sectioning. 

Figure 49 is a schematic diagram of S/N 132, which identifies 

these two sections and references these sections to the bolt 

hole letters and numbered acoustic cavities. Stereo optical 

microscopy was also performed on the two thruster sections 

prior to further sectioning. It was found that relief radius 

cracking extended around about half of the smaller section, 

from about AC number 10 through to the edge of AC number 14. The larger section was cracked nearly all the way 

around its circumference, extending from the edge of AC 

number 19 to 42. It should be noted that, based on the NESC 

analysis of the S/N 132 thruster pieces, AC numbers 14 and 16 

were incorrectly identified in the failure analysis in reference 48. It is more likely that the piece labeled AC number 14 

actually contained half of bolt hole A and incomplete portions 

of AC numbers 15 and 16; and the piece labeled AC number 16 actually contained bolt hole B and part of AC number 18. The corrected locations of these thruster pieces are 

shown by the cross-hatched regions in figure 49. 

The NDE results that were obtained during various points 

in the history of S/N 132 are also indicated in figure 49. The 

UT inspections that were performed in the 1979−1980 time 

frame by Marquardt and in 2004 by Boeing HB are given by 

the red and green lines, respectively, in figure 49. The 2004 

dye penetrant inspection results are indicated by the blue lines 

in the figure. There is good general agreement between the 

2004 UT inspections and the 2004 dye penetrant inspections in 

terms of relief radius cracking locations; and these indications 

agreed well with the observations made by stereo optical microscopy during the destructive examination of the actual 

hardware. The 1979−1980 UT inspections, however, did not 

adequately represent the extent or location of cracking in S/N 

132, which provides further evidence that the UT methods 

employed were neither reproducible nor reliable. 

A series of 10° sections (AA through RR) were cut subsequently by wire electrodischarge machining (EDM). The sections were cut from the injector flange, inward through the 

acoustic cavity wall, and across the injector face to the far 

edge of the acoustic cavity. A perspective of this sectioning is 

illustrated in figure 50. The remainder of S/N 132 is shown in 

figure 51 after all EDM was completed; some of the relief 

radius was intentionally left intact with the injector face to 

facilitate development of NDE methods. Scanning electron 

microscopy was used to examine the surface of the relief 

radius cracks in the 10° cut sections. Figure 52 shows a montage of images taken from section GG from S/N 132. The 

cracking is clearly evident with multiple initiations frequently 

observed along the relief radius surface. Sets of arrows in the 

figure show examples where the cracks run parallel to each 

other for some distance; apparently the cracks do not shield 

each other under the stress states imposed. If polished cross 

sections were obtained at the locations indicated by the 
NASA/TP—2005-214053 17 

arrows, then multiple cracks would be viewed in those polished cross sections. Section GG exhibited the typical degree 

of tightness of the crack at the surface. Figure 53 shows a 

range of tightness at the crack mouth in two different sections 

of the relief radius. In section QQ (fig. 53(a)), the crack is so 

tight that it is difficult to discern even by SEM. In contrast, 

section JJ contained a relatively wide open crack with little 

apparent branching; cracking was confined to a ledge in the 

relief radius for this section, and higher magnifications clearly 

reveal the intergranular nature of the cracking. The ledge of 

the relief radius is well above the root of the relief radius, part 

way up the injector flange (figs. 54(a) and (b)). 

3.2.3 Polished Section Observations 

Each EDM section had at least one side that was mounted 

and polished metallographically to obtain crack depth data 

prior to physically opening the crack. Figure 54(b) shows an 

example of one of these polished sections that contains a relief 

radius crack. The crack follows the grain boundaries, and thus 

some of the grains were cracked all the way around their 

periphery, which resulted in grains falling out during metallographic polishing. Hence, figure 54(b) shows typical polishing artifacts of pulled-out grains; the voids from these 

pulled-out grains do not represent actual crack widths. The 

micrograph also depicts the crack initiating at the relief radius 

surface at the previously mentioned ledge, which can be envisioned as a location where contaminants may become easily 

entrapped. Figure 55 shows section KK in an area where multiple cracks initiated at the relief radius; in this case, cracks 

with multiple branches initiated just above the relief radius 

ledge and another initiated at the base of the relief radius root. 

Again, grain pullout artifacts are observed. Between the tips of 

the parallel cracks in figure 55, a branch of one of the cracks, 

or a portion of another crack, is observed to extend well beyond the other two cracks. It is evident in figure 56 that the 

crack was intergranular along the entire depth and no change 

in fracture mode was observed in the polished sections. 

Several measurements were obtained from the relief radius 

cracks in the polished sections; these are depicted on the metallographic mount in figure 57. Actual crack depths were measured from the point of initiation along a straight line to the crack 

tip. The angle of inclination was determined from the counterclockwise angle between the outer surface of the AC wall and 

the straight line representing the actual crack depth. Finally, 

the projected crack depth was measured from the point of 

crack initiation to a line that begins at the crack tip and runs 

parallel to the outer surface of the AC wall. Table V summarizes the measurements that were taken from the polished 

surfaces of S/N 132. The table identifies each section by letter 

and its location on the circumference of the relief radius, based 

on the number of degrees (clockwise) from AC number 1. 

Three sections were polished on both sides. There is an absence 

of data between 0° and 50° and between 107° and 155° because 

those parts of the injector were not available for analysis. 

Table V indicates that most of the polished sections contained relief radius cracks. The percentage of the total distance 

of the crack to the acoustic cavity was also calculated for the 

projected crack depths in sections that contained acoustic 

cavities; some sections contained no acoustic cavities in the 

polished section. Table V shows that most of the projected 

crack depths measured were greater than 50 percent of the 

distance to the acoustic cavity; thus the crack depths were 

significant in S/N 132. Remaining ligament distances to the 

acoustic cavities are also provided in the table; it may be seen 

from these data that no cracks in the polished sections were 

observed to extend all the way to the acoustic cavities, 

although the cracks in sections OO and PP came within 

0.05 mm (0.002 in.) of the distance to the acoustic cavities. 

Table V also indicates that the angles of the cracks were very 

consistent, typically between 130° and 140°. The actual crack 

depths and projected crack depths are displayed graphically in 

figure 58, where the crack measurement is shown versus the 

location within the thruster, based on the number of degrees 

clockwise from AC number 1. The dashed line in the figure 

represents the projected distance from the relief radius to the 

acoustic cavity. Figure 58 shows that cracks occur over 

75 percent of the injector circumference, with the deepest 

cracking present in sections FF, GG, and NN. In these three 

sections, the crack front actually tunneled between two 

neighboring acoustic cavities, resulting in projected crack 

depths that were greater than the distance to the acoustic cavities. No pattern between the depth of cracking and location 

could be ascertained. 

3.2.4 Measurements After Crack Opening 

The polished sections were removed from the metallographic mounts, and the intergranular cracks were physically 

opened in the laboratory. Two-thirds of the sections were 

opened by placing them in a vise and hitting them with a hammer on the underside of the injector flange. Repeated hammer 

hits were made until the crack was opened more than 180°; at 

that point, vise grips were used in a limited reverse bending 

action to fracture the remaining ligament of material ahead of 

the crack. Some of the cracks were carefully back cut with a 

dry saw to facilitate the subsequent hammer openings. Onethird of the sections were placed in a vise with a large diameter dowel in front of the relief radius; the dowel and two edges 

of the sample enabled the crack to be opened more gently in 

the vise, using a three-point bend-type arrangement. It should 

be noted that all S/N 132 samples exhibited fully ductile overloads ahead of the intergranular crack tip, regardless of the 

crack opening procedure. 
NASA/TP—2005-214053 18 

The opened fracture surfaces of the intergranular cracks 

were examined extensively by SEM. The observations confirmed that the fracture surface was intergranular with no 

change in crack mode observed from the mouth to the tip of 

the crack (fig. 59). When multiple cracks were observed initiating on the relief radius, these cracks did not always connect 

subsurface to form a single crack front, as evident in figure 60. 

It should also be noted that, although the majority of the 

cracking had initiated in the relief radius, cracks were occasionally observed elsewhere in S/N 132. For example, shallow 

IGC was observed on the injector flange well outside of the 

relief radius. This was observed in polished mounts of section 

AA (figs. 61(a) and (b)) and after the lab opening attempt of 

section RR (figs. 61(c) and (d)). In both cases, the intergranular cracks extended only a few grains in from the flange surface. Cracks in this location would be consistent with the 

preweld etchant becoming entrapped after entering through the 

bolt holes, so it is not surprising to find affected thruster material in these areas. 

After crack openings were performed, it was found that 

relief radius cracks penetrated four acoustic cavities: AC 

numbers 19 and 20 (both in section GG), AC number 35 (section NN), and AC number 39 (section OO). An example of 

this is seen in figure 62(a), where the crack front is seen to 

extend into AC number 39; the crack tip fell just short of AC 

number 40. The micrograph in figure 62(b) also illustrates 

how the crack tunneled between two neighboring acoustic 

cavities; a more pronounced example of this is depicted in 

figure 63, where the crack front is clearly at a maximum depth 

between the acoustic cavities. This crack in section FF never 

penetrated either acoustic cavity. The examination of cracks of 

various depths suggests that the crack in this region was initially uniform in depth. As the crack front approached the 

acoustic cavities, the crack tip driving force in the region of 

the acoustic cavities became reduced relative to the crack tip 

region between the acoustic cavities, which resulted in the 

development of a tunneled crack front. The most likely cause 

for the tunneled crack front is the introduction of compressive 

residual stresses about the acoustic cavities due to the machining operation. It should be noted that no residual stress measurements were made and that the proposal of residual stresses 

in the region of the acoustic cavities is based solely on the 

observation of the shape of the crack front from the examination of S/N 132. 

Several measurements were made on the fracture surfaces. 

The fractures were oriented in two ways: with the viewing 

axis perpendicular to the crack face, which enabled actual 

crack depths to be measured; and with the viewing axis parallel to the acoustic cavity wall, which enabled projected crack 

depths and minimum ligament distances to be measured. In 

addition, ligament distances were obtained at the shortest 

distance from the crack front to the nearest acoustic cavity, as 

well as from the maximum crack depth to the nearest acoustic 

cavity; figure 64 illustrates that these minimum ligament 

distances were usually not the same. Table VI summarizes the 

measurements that were taken from the opened crack surfaces 

of S/N 132. The table identifies each section by letter and its 

location on the circumference of the relief radius, based on the 

number of degrees clockwise from AC number 1. Again, there 

is an absence of data between 0 and 50° and between 107° and 

155°, because those parts of the injector were not available for 

analysis. 

Table VI shows that only two sections, AA and BB, did not 

contain cracks; section QQ contained a very shallow crack that 

could not be fully opened even with careful back cutting. The 

remaining sections exhibited deep IGC. The column containing the minimum ligament distance to the acoustic cavities 

shows that three sections had cracks that penetrated all the 

way to the acoustic cavities, as indicated by the zero in the 

column. The crack profiles for the opened crack faces are 

depicted in figure 65; the crack measurements for maximum 

actual crack depth and maximum projected crack depth are 

shown versus the location within the thruster, based on the 

number of degrees from AC number 1. The dashed line in the 

figure represents the projected distance from the relief radius 

to the acoustic cavity. Figure 65 shows that cracking occurs in 

over 75 percent of the injector circumference, with the deepest 

cracking present in sections FF, GG, JJ, and NN. The crack 

fronts in these sections tunneled between two neighboring 

acoustic cavities and thus had projected crack depths greater 

than the distance to the acoustic cavities. Again no pattern 

between the depth of cracking and location could be ascertained. The cracking in S/N 132 was considerably deeper and 

more extensive than that observed in S/N 120, based on comparisons of figures 39 and 65. It is not known how the crack 

profiles of these two thrusters compare with the remainder of 

the thruster fleet. 

3.2.5 Observations for Crack Growth in Thrusters 

The fracture surfaces in S/N 132 were also examined carefully for any evidence of crack growth. Popcorn oxidation was 

observed along the crack front from the crack mouth to the 

crack tip, as seen in figures 66(a) to (c). No significant 

changes in the popcorn appearance were observed across the 

crack surface. However, regions of clean-looking grains with 

no popcorn were observed at the crack tip next to grains that 

were covered with popcorn, and an example of this is shown 

in figure 66(d). This clean appearance of some of the grains at 

the crack tip has been reported previously for S/N 132 in reference 48. Intergranular crack surfaces with no popcorn or 

other reaction products have also been produced by monotonic 

laboratory testing; this was shown previously in the C-ring 

specimen in figure 25(b) and in the uniaxial tensile specimen 

in figure 33. In contrast to S/N 132, the clean grains in the 

laboratory-tested specimens were observed across wide 
NASA/TP—2005-214053 19 

regions of the fracture and were not confined to the crack tips. 

EDS spectra from the grains without popcorn have reproducibly shown that they are minimally oxidized (fig. 25(c)). 

Several explanations have been proposed for the observation of clean grains at the crack tip in S/N 132. One explanation is that the crack tip is so tight that it limits oxidation at the 

crack tip. Although this may be feasible, the lack of bare 

grains at the crack tip in S/N 120 is hard to rationalize, 

because it would also be expected to have tight crack tips. The 

fact that minimally oxidized grains were observed in laboratory tests designed to reproduce the intergranular thruster 

failures keeps this tight crack tip theory as a strong contender. 

However, another possibility is that the intergranular cracks in 

S/N 132 grew in these very localized regions after thruster 

bakeout. It will be shown in section 3.3 that popcorn features 

develop only in the presence of fluorine and at temperatures 

greater than 204 to 232 °C (400 to 450 °F). If the cracks 

propagated during qualification testing in S/N 132, the lack of 

popcorn at the crack tip could be explained because the 

thruster temperatures during normal operations are significantly lower in magnitude and duration compared to those 

during bakeout, and thus popcorn would not be expected to 

form in those regions. Lattice dilation caused by the stress 

field at crack tips can enhance hydrogen permeation into these 

regions (ref. 49); the trapped hydrogen would be available to 

cause cracking under service loading at lower service temperatures, where materials are often more susceptible to hydrogen 

embrittlement. This cracking would be expected to have a 

limited extent since the original source of the hydrogen would 

no longer be present. However, the lack of bare grains at the 

crack tip in S/N 120 is again hard to rationalize because S/N 

120 was the fleet leader with 29 flights and over 12 000 firings 

(ref. 50), although S/N 132, as a qualification test article, 

underwent 100 simulated missions and 26 000 firings (refs. 33 

and 51). The final possibility is that the bare tip regions signify minimal crack growth during the unstable firing that 

occurred in S/N 132 during qualification testing. The short 

duration instability to a peak temperature of 304 °C (580 °F) 

would not have provided adequate time for significant oxidation and popcorn formation to occur. The service history and 

timing of the leak checks in S/N 132 are not inconsistent with 

this possible explanation. 

The three possible explanations given above for the observation of minimally oxidized grains at the crack tip in S/N 132 

cannot be ruled out. If the clean grains are interpreted as evidence for crack growth following manufacturing, the amount 

of crack growth would be extremely limited. Therefore, it has 

been concluded that no evidence was found in S/N 132 for 

appreciable crack propagation after manufacturing, with 

appreciable being defined as greater than 0.08 mm (0.003 in.). 

The lack of widely distributed clean grains at the crack tip is 

demonstrated in figure 67. This figure is a montage of BSE 

micrographs showing a typical distribution of minimally 

oxidized grains at a crack tip in S/N 132. It is clear that the 

clean-looking grains are not extensively distributed and are 

present only in isolated pockets along the crack front; these 

regions of minimal oxidation have been circled in the montage. 

A few other features were noted on the fracture surfaces of 

S/N 132. Figure 68 shows what appears to be slip planes on 

some of the minimally oxidized grain facets at the crack tip. 

This may provide evidence for a hydrogen embrittlement 

mechanism, which involves sweeping hydrogen into the plastic zone through dislocation transport from the crack tip. Hydrogen becomes trapped on the dislocations at the crack tip 

and preferential slip occurs locally. The uptake of hydrogen at 

the crack tip by this mechanism has produced intergranular 

fracture modes in the literature (refs. 52 and 53). Other regions 

of the oxidized fracture surface show what appear to be striated patterns of deposits on the grain facets (fig. 69). This has 

been observed previously in S/N 132 (ref. 48) and in counter 

bore cracks of S/N 120 (ref. 54). The reason for this striated 

pattern is unknown at this time; it is possible that this feature 

may be the result of popcorn formation along the slip planes 

observed in figure 68. 

3.2.6 Conclusions and Summary of Observations for 

S/N 132 

No evidence was found for appreciable crack propagation 

after manufacturing. Appreciable is defined to be greater than 

0.08 mm (0.003 in.) and corresponds to the longest lengths at 

the crack tips where “clean” grains were observed. 

(1) Cracking was intergranular from the mouth to the tip 

with no change in crack mode observed. 

(2) The intergranular crack faces contained significant oxidation with popcorn evident from the mouth to the tip, except 

for small pockets of clean, popcorn-free grains at the crack tip. 

These clean-looking grains were not observed in S/N 120. 

(3) Extensive cracking was observed in over 75 percent of 

the injector circumference. Cracking was significantly deeper 

and more extensive in S/N 132 compared to that in S/N 120. 

(4) Four acoustic cavities in S/N 132 were breached by 

IGC. 

3.3 Oxidation Tests to Reproduce Popcorn Oxidation 

3.3.1 Testing and Observations 

Additional tests were conducted to determine if the popcorn 

oxidation observed on the intergranular fracture surfaces of 

the thruster hardware could be reproduced. This testing would 

assist in determining if the intergranular cracks seen in thruster 

hardware were the result of manufacturing steps alone or if the 

thruster cracks were likely to have grown in service. An 
NASA/TP—2005-214053 20 

existing intergranular surface from previous laboratory testing 

on C−103 was selected as a specimen for subsequent oxidation 

testing, because it was a fracture surface that did not contain 

the numerous reaction products described in sections 3.1 and 

3.2. This specimen, shown previously in figure 25, was from 

C-ring testing and was only minimally oxidized based on 

detailed electron microprobe and EDS analyses. This opened 

intergranular fracture surface was then exposed isothermally 

in air at 316 °C (600 °F) for durations of up to 100 hr. The 

100-hr duration was selected as an upper limit, because thrusters undergo insulation bakeouts at 316 °C (600 °F) for 48 hr in 

air, and both S/Ns 120 and 132 have each undergone two 

48-hr bakeouts in their processing history. 

SEM and EDS were performed after each oxidation exposure of 5, 10, 20, 30, 50, and 100 hr at 316 °C (600 °F) to 

monitor the intergranular surface for possible changes in surface features and chemistry. After each oxidation interval, 

numerous grain facets on the fracture surface were examined 

for possible popcorn formation, since grain orientation has 

been shown to influence subsequent oxidation (ref. 44). Figure 

70 shows the same set of grains at a triple point as a function 

of exposure time at 600 °F. It is evident in the figure that no 

popcorn or cuboidal oxide formation was observed even after 

100 hr at 316 °C (600 °F), which is roughly equivalent to the 

exposure conditions from two thruster bakeout cycles. Figure 71(a) shows the EDS scan for the starting fracture surface, 

which is representative of minimally oxidized C−103 material. 

EDS scans obtained after 10, 50, and 100 hr at 316 °C (600 °F) 

are displayed in figures 71(b) to (d), respectively, and show 

that the biggest change in the fracture surface chemistry is 

observed after 10 hr, where the oxygen peak increased significantly in magnitude. Thus, the sample was oxidizing during 

the 316 °C (600 °F) exposures, but no other significant 

changes were evident. The results show that exposure times 

and temperatures consistent with a double bakeout cycle 

enable an oxide scale to form on C−103 but are insufficient for 

popcorn formation. It should be pointed out that the thruster is 

exposed to the out-gassing of water and carbon dioxide 

(ref. 55) from the thruster insulation material during the manufacturing bakeout, and this was not being simulated in these 

laboratory oxidation tests. 

Air exposures were also conducted at 399 °C (750 °F) for 

5 hr on the intergranular fracture surface after the 100-hr 

exposure at 316 °C (600 °F) and on virgin C−103 specimens 

with polished surfaces. This exposure temperature was selected to simulate the effects of an inadvertent temperature 

excursion during the manufacturing bakeout cycle and for 

comparison to oxidation experiments (ref. 44) on pure niobium. No popcorn oxidation formed on either the intergranular 

surfaces or the polished surfaces of C−103 after 5 hr at 399 °C 

(750 °F), which is inconsistent with the blistering observed 

after similar exposures in the literature were conducted on 

pure niobium (ref. 44). The reason for this difference in 

behavior is not known, although the alloying elements may 

have affected the oxidation rates and the breakaway oxidation 

phenomenon. Since the popcorn features on the thruster hardware surfaces were analyzed and were found to contain fluorine, it is believed that the absence of residual etchant on the 

C−103 test specimens contributes to the lack of popcorn formation. This is supported by the next set of tests performed on 

C−103 machined specimens and is described below. 

Two polished samples of C−103 were exposed to the 

preweld etchant consisting of HF, HNO3, and H2O in equal 

proportions. While wet, the two samples were laid one on top 

of the other, forming a “sandwich” specimen with the etchant 

entrapped. The sandwich was held together with stainless steel 

wire and dried overnight before an air exposure at 316 °C 

(600 °F) for 24 hr. The specimen was separated subsequently 

for examination. Numerous reaction products with different 

textures and colors were observed with some locations displaying popcorn-like oxidation, as seen in figure 72. The popcorn oxidation was observed in the regions that contained 

dried HF deposits. Similar tests were conducted with C−103 

sandwich specimens at 204, 399, and 816 °C (400, 750, and 

1500 °F) for 24 hr in air. The specimens exposed to 204 °C 

(400 °F) produced oxidized surfaces with small bumps present 

on the oxide scale; the bumps were similar in size to the popcorn oxidation observed on the thruster hardware, although the 

bumps did not have the cracked appearance that the popcorn 

typically exhibited. The bumps appeared to be precursors to 

the popcorn oxidation, and varied in density across the specimen surfaces. In contrast, the 399 and 816 °C (750 and 

1500 °F) exposures produced cracked oxide scales, but neither 

exposure produced popcorn oxidation. These tests suggest that 

a narrow range of temperature and the presence of fluorine are 

required to form popcorn oxidation on C−103. Specifically, 

for the cracked popcorn to form, it appears that a temperature 

greater than 204 °C (400 °F) and less than 399 °C (750 °F) are 

required. These data support the conclusion that the formation 

of popcorn oxidation on the intergranular cracks in the thruster 

hardware most likely occurred during the 316 °C (600 °F) 

manufacturing bakeout cycle. 

3.3.2 Conclusions and Summary of Oxidation Testing 

The intergranular cracks seen in the thruster hardware were 

the result of manufacturing steps alone, since popcorn features 

observed on the hardware IGC could only be reproduced with 

316 °C (600 °F) manufacturing bakeout cycles in the presence 

of HF-containing etchants. 

(1) A narrow range of temperature and the presence of 

fluorine are required to form popcorn oxidation with a cracked 

appearance on C−103. 

(2) Normal thruster operation is not expected to form 

cracked popcorn oxidation because C−103 specimens exposed 
NASA/TP—2005-214053 21 

to 204 °C (400 °F)/24 hr only produced oxidized surfaces with 

small bumps that appeared to be precursors to the popcorn 

oxidation. 

(3) Manufacturing bakeout cycles conducted without HFcontaining etchants did not produce popcorn oxidation after 

100 hr of exposure. 

3.4 Fracture Surface Analysis by Energy Spectroscopy for Chemical Analysis (ESCA) 

3.4.1 Oxide Analysis 

ESCA is a surface analysis technique that was used to identify elements and their oxidation state that were present on the 

fracture surface of intergranular cracks in S/Ns 120 and 132. 

Organic and inorganic compounds can be identified with 

ESCA by comparison of binding energy spectra from known 

compounds. Generally, the best quantitative results can be 

obtained with polished specimens, but the presence of elements and qualitative trends can be established on fracture 

surfaces. In addition, depth profiling through the surface layers can be performed with ion-milling; and again, although 

this can be done on a fracture surface, this technique is best 

accomplished on a flat specimen where a uniform amount of 

material removal can be accomplished. 

ESCA was performed first on the surface of the relief radius 

crack near hole L of S/N 120. The original intent was to identify reaction products and determine their distributions across 

the intergranular fracture surface by taking ESCA scans at 

multiple locations. Figure 73 shows the fracture surface of this 

relief radius crack and locations A through F where ESCA 

analysis was performed. The dotted line superimposed on the 

image of the fracture surface indicates the approximate locations for the crack tip. Location C represents the lab overload 

area, whereas the other locations are from the intergranular 

crack surface. The extremities of the crack surface, as well as 

areas having different surface colorations, were intentionally 

selected for analysis, in an effort to determine if the spectra 

differed as a function of location, thereby indicating possible 

progression of the crack over time. 

Figure 74(a) shows the ESCA spectra obtained from the lab 

overload area at location C; it is evident that four peaks are 

resolved for the spectra. Typical ESCA spectra from metallic 

niobium surfaces would be expected to show a binding energy 

doublet, or two peaks of equal magnitude, for Nb 3d5/2 and 

3d3/2 at 202.0 and 204.7 eV, respectively (ref. 56). As the 

niobium surface becomes oxidized, the peaks in the spectra 

shift to higher binding energy values, with the higher oxides 

(e.g., Nb2O5) shifting more than the lower oxides (e.g., NbO). 

Comparison of the binding energy line positions for the peaks 

in figure 74(a) to reference spectra in the literature (refs. 56 

and 57) for niobium, NbO, NbO2, and Nb2O5 indicates a close 

match to niobium metal and Nb2O5. The two rightmost peaks 

in figure 74(a) represent the niobium doublet and the two 

leftmost peaks represent Nb2O5. Thus, based on these data, the 

lab overload area has a slightly oxidized niobium surface. 

In contrast, the typical spectra obtained from the intergranular crack surfaces at five locations (A and B and D to F) are 

shown in figure 74(b), where only three peaks are resolved for 

those spectra. Comparison of the line positions for the peaks in 

figure 74(b) to the reference spectra in the literature (refs. 56 

and 57) reveals a close match to NbO and Nb2O5; no metallic 

niobium doublet is evident. The rightmost peak in the figure 

represents NbO, the leftmost peak represents Nb2O5, and the 

middle peak has the highest magnitude from the overlap of 

additional peaks from NbO and Nb2O5. The spectra in figure 74(b) are typical for more heavily oxidized surfaces that 

contain a mixture of oxides (ref. 56). Thus, the intergranular 

crack surface has a thicker and more complex oxide scale than 

that of the lab overload region. Table VII lists the line positions for the low oxide (NbO) Nb3d5/2 and for the high oxide 

(Nb2O5) Nb3d3/2 obtained from locations A and B and D to F 

in S/N 120. It is evident that the line positions were similar 

across the entire intergranular crack surface from mouth to tip. 

Additionally, the areas under the peaks for the low and high 

oxides were integrated, and the resulting percentages of total 

volume for each oxide on the analyzed surface are also given 

in table VII. These data indicate that the low oxide, NbO, is 

consistently the majority phase in the oxide scale of the intergranular crack surface. Thus, the same oxides were found in 

approximately the same volume fraction across the fracture 

surface, which provides further evidence that the intergranular 

thruster cracks did not form progressively over time. 

3.4.2 Analysis of Other Compounds 

The ESCA spectra from the intergranular crack surfaces 

revealed the presence of additional peaks that were associated 

neither with the niobium substrate nor with the above mentioned oxides. Figure 75(a) shows spectra containing two peaks 

for fluorine and for oxygen and multiple small peaks for niobium, and figure 75(b) shows spectra at lower binding energies 

where three peaks for carbon and for niobium are seen. The 

rightmost peaks of fluorine and oxygen in figure 75(a) are 

believed to be associated with NbO2F, a reaction product 

described earlier as resulting from a chemical reaction between 

NbF5, oxygen, and residual moisture in the air (ref. 32). The 

rightmost carbon peak that is barely above background levels in 

figure 75(b) is likely the result of contamination from handling. 

The niobium peaks in figure 75(b) were resolved in figure 74(b) 

and were described in the previous section. 

The remaining peaks for fluorine, oxygen, and carbon in 

figure 75 were relatively large in magnitude and were found 

consistently over the analyzed fracture surface of S/N 120; 

these peaks were not associated with a known reaction product 

and thus were explored in more detail. A subsequent review of 
NASA/TP—2005-214053 22 

the injector processing steps revealed that a perfluoropolyether 

(PFPE) lubricant, Krytox 143AC, may have been applied to 

the bolts and/or bolt holes on the flange of the niobium injector during thruster assembly (ref. 55). Krytox 143AC was also 

used to lubricate the O-ring on the flow test fixture during four 

separate flow tests that were performed after various machining steps during manufacturing (ref. 55). Krytox material 

contains fluorine, oxygen, and carbon, and as a result of the 

unknown fluorine, oxygen, and carbon peaks obtained from 

the thruster fracture surface, an ESCA analysis was performed 

directly on samples of this PFPE lubricant. A comparison 

between the line positions obtained from Krytox 143AC and 

the additional (leftmost) fluorine, oxygen, and carbon lines on 

the thruster hardware surfaces indicated an excellent match. 

Furthermore, it should be noted that the PFPE lubricant 

migrated at least 1 mm away from its original liquid meniscus 

during ESCA analysis, and thus it was determined that the 

material has a low surface tension. Therefore, it is not hard to 

imagine that the Krytox material may have also migrated into 

the relief radius, and possibly into a preexisting crack, after 

being applied to the bolt hole regions of the injector. 

Ion-milling was also performed at locations A and B and D 

to F (fig. 73) to remove surface layers in order to determine 

depth profiling trends. The maximum sputtering duration was 

300 s, which removed about 100 nm of material, and binding 

energy spectra were obtained during the sputtering operation. 

It was found that as material was being removed from the 

fracture surface during sputtering, the peaks associated with 

the Krytox lubricant were reduced in magnitude, and the peaks 

corresponding to NbO2F became more prominent. Carbon was 

not detected in the subsurface layers corresponding to a depth 

of 100 nm. These results are consistent with the trends determined from Auger depth profiling (ref. 4) and from the x-ray 

line scans presented in figure 47. Eventually, as more of the 

fracture surface material was sputtered away with ion-milling, 

the resulting ESCA spectra only contained peaks representing 

niobium oxides. The similarity in the ESCA results obtained 

before and after ion-milling at various locations on the S/N 

120 fracture surface suggests that the IGC was not progressive, which supports the hypothesis that the IGC occurred 

during manufacturing. 

The fracture surface from a relief radius crack near bolt 

hole B in S/N 132 was also examined by ESCA. This was 

from a deeper crack than that analyzed in S/N 120, and the 

location of this cracked section within S/N 132 is illustrated in 

figure 49. It should be noted that reference 48 incorrectly 

identified this section as being adjacent to acoustic cavity 

number 16, when it was actually from the cross-hatched section in figure 49 between acoustic cavity numbers 17 and 18. 

The fracture surface analyzed from S/N 132 is shown in figure 76, and the specific areas that were analyzed by ESCA are 

indicated in the figure at locations A to G. The fracture surface 

is entirely intergranular, and no lab overload region is present 

in figure 76. Although the entire fracture surface examined in 

S/N 132 is oxidized, the spectra from the ESCA analysis 

indicate lower peak intensities for oxygen (and hence lower 

amounts of oxide) at the crack tip near locations C, D, and F, 

in comparison to those obtained from the crack tip in S/N 120. 

This result is consistent with the fractography and EDS 

results, which indicate some minimally oxidized areas in 

localized regions at the crack tip, previously shown in figures 66(d) and 67. Based on the ESCA spectra line positions, 

the types of oxides and oxy-fluorides present on the crack 

surface S/N 132 match those present on the intergranular crack 

surface of hole L in S/N 120. Evidence for the Krytox lubricant was also found based on additional peaks for fluorine and 

carbon in the ESCA spectra, although the lower intensity of 

these peaks indicates that less Krytox was present on the S/N 

132 fracture surface compared to that of S/N 120. Moreover, 

ESCA results also indicate even lower quantities of the lubricant were present in S/N 132 at the crack tip locations D and F 

compared to the remainder of the S/N 132 fracture surface. 

The presence of the PFPE lubricant on the crack surfaces of 

two thrusters is considered noteworthy, because Krytox 

143AC has been shown to break down by reacting with surface oxides to produce metal fluorides (ref. 45). The metal 

fluorides are strong Lewis acids that further decompose the 

PFPE lubricant in an autocatalytic process; the smaller segments of the polymer chain react with moisture, and trace 

amounts of HF are released. Thermogravimetric analyses were 

performed and coupled to gas cell in a fourier transform infrared spectrometer (ref. 46), in order to confirm the expected 

reactions between Krytox 143AC and C−103 in the presence 

of moisture in an air environment. It was determined that the 

Krytox material unzips into smaller segment oligomers at 

temperatures close to the thruster bakeout temperature 

(ref. 46), and trace amounts of HF were detected after 23 min 

at 399 °C (750 °F). A similar compound, Fomblin Y, with an 

identical ESCA signature was also analyzed and was found to 

breakdown in an air environment at 299 °C (570 °F) in the 

presence of C−103 (ref. 46). The breakdown of either lubricant material would cause the release of small amounts of HF, 

which would have the potential of causing IGC in a manner 

similar to that which occurred by the preweld etchant during 

316 °C (600 °F) bakeout. Testing of C−103 specimens produced IGC and surface features consistent with that observed 

in service after exposure to Krytox and heating to 316 °C 

(600 °F) for 72 hr (ref. 42). However, based on the ESCA 

analyses performed during ion-milling, the Krytox material 

appears to be located only in the top layers of the intergranular 

crack surface, suggesting that the Krytox may have been 

introduced after the crack formed and the crack surface 

became oxidized. It should also be noted that, based on the 

ESCA alone, it cannot be determined if any residual PFPE 

lubricant on the thruster hardware is still capable of being 

reactive upon exposure to elevated temperatures, because the 

breakdown products of the PFPE lubricant are expected to 

give similar signatures to the lubricant itself. 
NASA/TP—2005-214053 23 

3.4.3 Summary and Conclusions From ESCA Analyses 

The intergranular cracks did not form progressively over 

time. 

(1) The same oxides (NbO and Nb2O5) were found across 

the fracture surface from the crack mouth to the crack tip in 

approximately the same volume fraction. 

(2) Evidence was found for Krytox lubricant in the top layers of the fracture surface. 

(3) Ion-milling showed similar depth profiling results 

through the thickness of the reaction products over the entire 

fracture surface: niobium oxides and Krytox were found on 

the surface; Krytox diminished as NbO2F became more 

prominent beneath the surface layers; and niobium oxides 

were present beneath NbO2F and above the parent material. 

(4) Krytox migrates due to its low surface tension. Since it 

was found only on the surface and not in layers adjacent to the 

parent material, Krytox may have migrated into a preexisting 

relief radius crack. 

Section 4: Exposures to Sodium Hydroxide via Oakite Rustripper Immersion 

4.1 Background 

During a chamber repair and rejuvenation process in March 

2005 at WSTF, it was discovered that a small number of injector assemblies had been erroneously immersed in Oakite Rustripper (Oakite Products, Inc., Berkeley Heights, NJ) during a 

cleaning operation. This process had been occurring on all 

thruster chamber repair and rejuvenations since 1999 (ref. 58). 

Oakite Rustripper contains a nominal concentration of 13 wt% 

sodium hydroxide (NaOH), an alkali that is well known to 

embrittle niobium and its alloys by the absorption of hydrogen. Niobium is attacked by hot (98 °C (208 °F)) NaOH in 

concentrations from 1 to 10 wt% (refs. 18 and 36) and is also 

attacked at room temperature by NaOH in concentrations 

ranging from 1 to 40 wt% (refs. 35 to 37). As part of a prior 

failure investigation in 1979 on injector S/N 128 (ref. 55), it 

was demonstrated that long exposures of Oakite Rustripper 

caused intergranular and transgranular failures at 79 °C 

(175 °F) in notched C−103 tensile specimens that were preloaded to 414 MPa (60 ksi). More recent testing (ref. 59) 

showed some loss of ductility and mixed mode fractures with 

4-hr exposures to Oakite Rustripper at 79 °C (175 °F), and 

severe embrittlement was observed after 20-hr exposures; 

however, no apparent degradation was observed during qualitative testing with 20 min rust stripper exposure times that 

approximated the process durations at WSTF. 

The concern over the Oakite Ruststripper exposures became 

heightened after the NESC Materials SPRT reevaluated the 

results from fractography that had been conducted earlier at 

GRC and Boeing HB on S/N 120 and at GRC, LaRC, and 

MSFC on S/N 132. S/N 120 had been exposed to Oakite Rustripper during its chamber repair and rejuvenation in 2004, and 

had exhibited embrittlement in the form of transgranular 

cleavage on the lab-induced fracture surface that was well 

ahead of the intergranular cracks and adjacent to acoustic 

cavities. This was consistently observed on the opened crack 

surfaces in S/N 120, irrespective of the laboratory in which the 

fracture was induced, although only a limited number of lab 

openings had been performed on that injector. It was postulated that the Oakite Rustripper would have come into direct 

contact with the acoustic cavities during the immersion process, thereby promoting hydrogen absorption and diffusion of 

hydrogen in the regions where the brittle behavior was 

observed. In contrast, 16 sections were opened on S/N 132 at 3 

different NASA centers, and 100 percent of these crack openings resulted in completely ductile features ahead of the intergranular crack and adjacent to the acoustic cavities. Injector 

S/N 132 was never exposed to Oakite Rustripper. Fracture 

behavior appeared at that time to be correlated with exposures 

to Oakite Rustripper. 

Although less than 10 injector assemblies have actually 

undergone repair and rejuvenation processes at WSTF, 2 such 

injectors had been installed on OV−103 for STS−114, and 

thus had been exposed to Oakite Rustripper immersions. Thus, 

the concern was raised by the NESC Materials SPRT that 

existing intergranular thruster cracks might propagate during 

service into the regions between the intergranular crack tips 

and the acoustic cavities, where elevated hydrogen contents 

may exist due to these erroneous NaOH immersions. The 

fractographic evidence that was provided by the NESC Materials SPRT showed the potential adverse effect of sodium 

hydroxide and contributed to a flight constraint on OV−103 

that was initiated by Safety and Mission Assurance. The fractographic evidence also provided the impetus for additional 

testing of thruster material. 

4.2 Testing With Oakite Rustripper 

An intensive testing program ensued at WSTF to address 

this concern regarding Oakite Rustripper immersions. Representatives from the Orbiter Materials and Processing Program 

and their contractors, Safety and Mission Assurance, and the 

NESC Materials SPRT participated directly in the mechanical 

testing and resultant fractography. Test specimens were 

machined from a fleet lead primary thruster injector in the 

following conditions: as-is with no Oakite Rustripper exposure; with 30-min Oakite Rustripper exposures performed at 

79 °C (175 °F) in the chemical laboratory at WSTF; and with 

20-min Oakite Rustripper exposures performed at 66 °C 

(150 °F) using equipment and procedures closely representing 

the nominal cleaning procedures at WSTF. Fleet leader 

specimens with and without Oakite Rustrippper exposures 
NASA/TP—2005-214053 24 

exhibited ductile rupture and similar tear resistances based on 

load versus displacement curves, which suggests that the 

WSTF Oakite Rustripper immersions did not allow sufficient 

hydrogen absorption to degrade mechanical behavior. Similar 

testing was also conducted on S/N 120, both in the 

as-processed condition and after a subsequent 316 °C 

(600 °F)/48-hr bakeout. The WSTF testing showed only a 

slight drop in the tear resistance of the S/N 120 material compared to that of the fleet leader specimens, although these 

small differences could have been due to testing variability 

with a non-standard specimen design. Fractography of the 

tested S/N 120 showed mixed-mode fracture and small areas 

of 100 percent transgranular cleavage, which suggests that S/N 

120 has been embrittled by elevated hydrogen levels. The 

source(s) of the apparent higher hydrogen content in S/N 120 

could not be conclusively established, but it appeared that the 

hydrogen had accumulated from a variety of potential sources, 

and not solely from the WSTF repair and rejuvenation process. Subsequent tests conducted by the NESC Materials SPRT 

at LaRC indicated only a very small decrease in the fracture 

toughness and no significant change in the fatigue crack 

growth rate when the S/N 120 specimens were compared to 

the fleet leader specimens. 

Therefore, it was concluded that the observed differences in 

behavior between S/N 120 and the fleet leader with and without Oakite Rustripper exposures were very small. As a result, 

the tests demonstrated that the C−103 material was not significantly affected by the 20-min Oakite Rustripper exposures at 

WSTF, and thus all thrusters given the Oakite Rustripper 

immersions during the chamber repair and rejuvenation were 

cleared for flight. However, since C−103 has the potential of 

being adversely affected by NaOH and by hydrogen uptake, it 

was recommended by the NESC Materials SPRT and Safety 

and Mission Assurance that the Oakite Rustripper immersion 

step be removed from the WSTF chamber repair and rejuvenation processing, in order to reduce the risk of additional 

hydrogen uptake in thrusters repaired in the future. As result 

of this recommendation, corrective action was taken, and an 

audit of the thruster repair and rejuvenation processes at 

WSTF was conducted. 

4.3 Summary and Conclusions From Oakite 

Rustripper Testing 

All thrusters given the Oakite Rustripper immersions during 

the chamber repair and rejuvenation were cleared for flight. 

Testing results indicated that the C−103 thruster material was 

not significantly affected by the 20-min Oakite Rustripper 

exposures at WSTF. 

(1) Only a slight drop in the tear resistance, a very small 

decrease in the fracture toughness, and no significant change 

in the fatigue crack growth rate were observed in the S/N 120 

material compared to that of the fleet leader specimens. 

(2) Fractography of the tested S/N 120 showed mixedmode fracture and small areas of 100 percent transgranular 

cleavage, which suggests that S/N 120 has been embrittled by 

elevated hydrogen levels but not sufficiently to degrade 

mechanical properties. 

(3) Recent testing, 1979 testing, and literature data demonstrated that niobium alloys can be degraded significantly by 

NaOH under certain conditions, particularly by longer 

exposures. 

(4) Due to the risk of additional hydrogen uptake in thrusters immersed in sodium hydroxide, a recommendation was 

made to remove Oakite Ruststripper immersions from the 

WSTF chamber repair and rejuvenation process. 

Conclusions 

1. RCS thruster intergranular cracking (IGC) is due to 

hydrogen embrittlement. 

 a. Hydrogen production on C−103 is thermodynamically 

assured in the presence of hydrofluoric-acidcontaining etchants. 

 b. C−103 has been shown to be embrittled when 

charged with hydrogen. 

 c. Dissolution processes (i.e., stress corrosion and hotsalt cracking) will not result in IGC for this C−103 

material. 

 d. Root cause identified in 1982 failure analyses was 

not substantiated. 

2. Thruster IGC was produced during manufacturing as a 

result of processing of thrusters in the presence of 

fluoride-containing acids. 

 a. Laboratory tests to simulate processing conditions 

have produced IGC with surface features consistent 

with that observed in service. These conditions 

included exposure to the niobium etchant 

(33%HF:33%HNO3:33%H2O) followed by a 316 °C 

(600 °F) exposure in air for 48 hr. 

 b. Oxide features (termed as “popcorn”) have been 

observed on the fracture surfaces. These popcorn features have only been replicated in the presence of 

fluorine and oxygen when heated to temperatures 

greater than 204 °C (400 °F), suggesting that the 

thruster fracture surfaces were exposed to the processing bakeout. 

 c. A fluorinated lubricant has been found on the fracture 

surfaces of serial numbers (S/Ns) 120 and 132. Krytox 143AC (DuPont, Wilmington, DE) is the fluorinated lubricant that is applied to the thrusters during 

processing. These lubricants have been shown to 

react with metal oxides and breakdown at thruster 

bakeout temperatures, thereby releasing small quantities of HF (refs. 45 and 46). New test results (ref. 42) 
NASA/TP—2005-214053 25 

produced IGC and surface features consistent with 

that observed in service after exposure to Krytox and 

heating to 316 °C (600 °F) for 72 hr. 

 d. Failure analyses of cracked thrusters have shown 

cracking to be 100 percent intergranular. No change 

in fracture mode has been observed at the crack tip. 

3. Appreciable environmental crack propagation does not 

occur after manufacturing. 

 a. Examination of the oxide features has shown similar 

features from the crack mouth to the crack tip, and 

has provided no indication of appreciable (greater 

than 0.08 mm (0.003 in.)) crack propagation. 

 b. The same oxides (NbO and Nb2O5) were found 

across the fracture surface from the crack mouth to 

the crack tip in approximately the same volume fraction. Depth profiling through the thickness of the 

reaction products on the IGC provided further confirmation that the reaction products did not change 

from the crack mouth to the crack tip. 

 c. Laboratory tests to determine environmental crack 

propagation conditions during service have not 

revealed any crack growth. Here, specimens are 

exposed to the processing conditions (exposure to the 

niobium etchant and baked at 316 °C (600 °F) for 

48 hr under stress) to produce IGC. The specimens 

are then exposed to a high humidity environment and 

monitored for crack propagation. 

4. The C−103 thruster material was not significantly 

affected by the 20-min Oakite Rustripper exposures during the thruster repair and rejuvenation processes at 

WSTF. All thrusters given the Oakite Rustripper immersions during the chamber repair and rejuvenation were 

cleared for flight. 

 a. Only a slight drop in the tear resistance, a very small 

decrease in the fracture toughness, and no significant 

change in the fatigue crack growth rate were 

observed in the S/N 120 material compared to that of 

the fleet leader specimens. 

 b. Fractography of the tested S/N 120 showed mixedmode fracture and small areas of 100 percent transgranular cleavage, which suggests that S/N 120 had 

been embrittled by elevated hydrogen levels but not 

sufficiently to degrade mechanical properties. 

 c. Recent testing, literature data, and testing from 1979 

all demonstrated that niobium alloys can be degraded 

significantly by NaOH under certain conditions, particularly by longer exposures. 

 d. Due to the risk of additional hydrogen uptake in 

thrusters immersed in sodium hydroxide, a recommendation was made to remove Oakite Ruststripper 

immersions from the WSTF chamber repair and rejuvenation process. 

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NASA/TP—2005-214053 26 

22. Grossbeck, M.L.; and Birnbaum, H.K.: Low Temperature Hydrogen Embrittlement of Niobium⎯Microscopic Observations. 

Acta Metall. vol. 25, no. 2, 1977, pp. 135−147. 

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Niobium⎯High Temperature Behavior. Acta Metall., vol. 26, 

no. 11, 1978, pp. 1781−1788. 

24. Spitzig, W.A.; Owen, C.V.; and Scott, T.E.: Effects of Interstitials and Hydrogen-Interstitial Interactions on Low Temperature 

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A, vol. 17A, no. 7, 1986, pp. 1179−1189. 

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nos. 1−2, 1989, pp. 339−366. 

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in Fe, Nb, and Pd. Scripta Metall., vol. 16, no. 4, 1982, 

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Alloy, Tantalum, and Ta-10W Alloy. Corrosion, 1991, 

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pp. 17−24. 

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FL, 2001, Niobium 2001 Ltd., Bridgeville, PA, 2002, 

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39. Subramanyan, P.K.: Comprehensive Treatise of Electrochemistry. Plenum Press, New York, NY, 1981, pp. 411−462. 

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45. Jones, William R., Jr.: The Properties of Perfluoropolyethers 

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Fomblin, Krytox, and Krytox With Niobium Alloy. April 19, 

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Cracks (CAR AE2746), Boeing HB Report, Nov. 2, 2004. 

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50. Fitzgerald, E.J.: RCS Primary Thruster Rationale for Flight. 

Presented at Orbiter Project Office Tagup, Johnson Space Center, June 1, 2005. 

51. Curtis, C.: United Space Alliance. Materials and Processes RCS 

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CA, Dec. 15, 2004. 

52. Shin, K.S.; Austin, N.J.; and Kim, S.S.: Effect of Hydrogen on 

Mechanical Behavior of a 2090 Al-Li Alloy. Hydrogen Effects 

on Material Behavior, Neville R. Moody and Anthony W. 

Thompson, eds., TMS, Warrendale, PA, 1990, pp. 1023−1031. 

53. Pressouyre, G.M.: Trap Theory of Hydrogen Embrittlement. 

Acta Metall., vol. 28, no. 7, 1980, pp. 895−911. 

54. Reid, M.: Boeing Rocketdyne, Nov. 2004 (unpublished results). 

55. Failure Investigation Intergranular Crack Causing Injector 

Flange Leakage⎯T.A. P/N 235216, S/N 128, F.M.R. 

No. 779−056, S−1494, MPM No. 18.343, July 1979 (The 

Marquardt Company Process Memorandum). 

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Phys. Chem. Sol., vol. 36, no. 6, 1975, pp. 485−491. 

57. Fontaine, R., et al.: Chemical ESCA Shift in Niobium Oxide 

Series. J. Electron Spectrosc. Relat. Phenom., vol. 10, no. 4, 

1977, pp. 349−357. 

58. Kelly, T.: Boeing WSTF. Presented to RCS Thruster Materials 

and Processing Team, March 24, 2005 (teleconference). 

59. Jacobs, J.: RCS Thruster Processing Escape: Oakite Rustripper. 

Presented to RCS Thruster Materials and Processing Team, 

April 21, 2005 (teleconference).

Glenn Research Center 

National Aeronautics and Space Administration 

Cleveland, Ohio, December, 2005 
NASA/TP—2005-214053 27 

TABLE I.—ORIGINAL SPECIFICATIONS FOR 

NIOBIUM INJECTOR MATERIAL 

Element Nominal composition Permitted variation 

Minimum Maximum (Under minimum or over maximum) 

Hafnium 9.0 wt% 11.0 wt% 0.15 wt% 

Titanium 0.7 wt% 1.3 wt% 0.07 wt% 

Zirconium 0.7 wt% 0.05 wt% 

Tungsten 0.5 wt% 0.03 wt% 

Tantalum 0.5 wt% 0.05 wt% 

Carbon 150 ppm None 

Oxygen 225 ppm None 

Nitrogen 150 ppm None 

Hydrogen 15 ppm None 

All others 3000 ppm None 

Niobium Balance 

Grain size of forgings shall be predominantly ASTM–E112 5 or finer with 

occasional grains as large as 3 permissible. 

Hardness (Rockwell B): 70–95

TABLE II.—OPEN CIRCUIT POTENTIALS (EOC) for C–103 AND Ti-6Al-4V 

EOC (V vs. SCE)

C–103 

EOC (V vs. SCE)

Ti-6Al-4V 

EOC (V vs. SCE) 

C–103/ Ti-6Al-4V 

couple 

Er* (V vs. SCE) 

Water reduction 

Dilute HF (pH ≈ 3.5) –0.478 –1.076 –1.010 –0.451 

1 wt% NaCl –0.383 –0.280 –0.355 –0.625 

*Er (water reduction) = –0.241 VSCE – (0.06 V × pH). 

TABLE III.—SCRATCH POTENTIALS (Escratch) AND CALCULATED HYDROGEN 

FUGACITY (fH2

) for C–103 IN THE DILUTE HF AND 1 wt% NaCl SOLUTIONS 

 Escratch (V vs. SCE)

C–103 

Er* (V vs. SCE)

Water reduction

fH2 (atm) 

C–103 scratch 

fH2 (atm) 

C–103 open circuit

Dilute HF (pH ≈ 3.5) –0.910 –0.451 2×1015 7.94 

1 wt% NaCl –1.285 –0.625 1×1022 8.58×10–9

*Er (water reduction) = –0.241 VSCE – (0.06 V × pH)
NASA/TP—2005-214053 28 

TABLE IV.—UNIAXIAL TENSION TESTING FOR ENVIRONMENTAL CRACKING 

Specimen Solution Exposure stress Bakeout Failure Microscopy

1 HF/H20/NaCl ≈ 198 MPa (28.8 ksi) 

–900 mVSCE

None Failure < 40.5 hr Nearly 100 percent cleavage

2 HF/H20/NaCl ≈ 177 MPa (25.6 ksi) 

–750 mVSCE

None Failure 26.5 hr Nearly 100 percent cleavage

3 HF/H20 ≈ 177 MPa (25.6 ksi) 

–750 mVSCE 24 hr

None Strained to 

failure

Embrittled zone visible 

No intergranular cracking or cleavage

4 HF/H20 ≈ 177 MPa (25.6 ksi) 

–600 mVSCE 48 hr

177 MPa (25.6 ksi) 

316 °C (600 °F) 48 hr

Strained to 

failure

Embrittled zone visible 

No intergranular cracking or cleavage

5 HF/H20 ≈ 39.8 MPa (5.76 ksi) 

–1 VSCE 90 hr

None Strained to 

failure 

Failed in grip 

section

Transgranular cleavage plus intergranular 

cracking

6 HF/H20 ≈ 39.8 MPa (5.76 ksi) 

–1 VSCE 90 hr

39.8 MPa (5.76 ksi) 

316 °C (600 °F) 17 hr

Strained to 

failure

Ductile rupture

7 HF/H20 No load –1 VSCE 90 hr None Strained to 

failure 

Failed in grip 

section

Nearly 100 percent cleavage

8 HF/H2O No load –1 VSCE 90 hr None Strained to 

failure 

Failed in grip 

section

Nearly 100 percent cleavage

9 HF/H2O No load –1 VSCE 90 hr 26.5 MPa (3.85 ksi) 

316 °C (600 °F) 17 hr 

Strained to 

failure 

Intergranular cracking up to 0.25 mm 

along outer surface 

10 HF/H20/NaCl ≈ 221 MPa (32.0 ksi) 

–100 mVSCE 91.5 hr

None Uniform attack No observed local attack
NASA/TP—2005-214053 29 

TABLE V.—SUMMARY OF CRACK MEASUREMENTS FROM POLISHED SURFACES 

S/N 132 

10° slice 

NASA center Degrees from 

AC number 1 

Projected 

crack depth, 

in. 

Percent of total 

distance to 

acoustic cavity 

(projected 

crack depth), 

percent 

Actual

crack 

depth, 

in. 

Ligament 

distance to 

acoustic cavity, 

in. 

Angle of crack 

(degrees CCW 

from acoustic 

cavity wall) 

AA GRC 50 0 0 0 0.195 No crack 

BB MSFC 60 0 0 0 No cavity No crack 

CC polished side LaRC 70 0 0 0 No cavity No crack 

CC opposite side LaRC 80 N/A >50 N/A No cavity ~135 

DD GRC 80 0.125 58 0.166 0.092 132 

EE MSFC 90 0.080 48 0.107 0.087 131 

FF remnanta LaRC 107 0.208 90 0.295 0.023 135 

GG GRC 155 0.216 94 0.286 0.015 133 

HH LaRC 175 0.057 N/A 0.065 No cavity 144 

II MSFC 205 0.139 80 0.208 0.035 138 

JJ GRC 225 0.145 72 0.190 0.055 131 

KK max. crack 

polished side 

LaRC 240 0.123 79 0.160 0.032 140 

KK opposite side LaRC 250 0.113 66 0.148 0.058 139 

LL MSFC 260 0.096 48 0.133 0.106 134 

MM LaRC 275 0.113 68 0.156 0.053 136 

NN GRC 295 0.181 68 0.236 0.085 133 

OO MSFC 325 0.166 99 0.228 0.002 133 

PP max. crack 

polished side 

LaRC 335 0.172 99 0.217 0.002 142 

PP opposite side LaRC 345 0.041 N/A 0.080 N/A 120 

QQ GRC 345 0.052 24 0.069 0.164 132 

RR remnanta GRC 358 0.014 8 0.016 0.161 116 a

Remnants are less than 10°. 
NASA/TP—2005-214053 30 

TABLE VI.—SUMMARY OF CRACK MEASUREMENTS FROM OPENED CRACK FACES 

S/N 132 

10° slice 

NASA center Approx. 

degrees from 

AC number 1 

for cut section 

Maximum 

projected 

crack depth, 

in. 

Maximum

actual crack 

depth, 

in. 

Minimum 

ligament 

distance to 

acoustic cavity, 

in. 

Ligament distance

to acoustic cavity 

at maximum 

crack depth, 

in. 

AA GRC 55 0 0 N/A N/A 

BB MSFC 65 0 0 N/A N/A 

CC LaRC 75 0.152 0.215 0.093 0.093 

DD GRC 85 0.105 0.115 0.071 0.071 

EE MSFC 95 0.151 0.199 0.016 0.016 

FF remnanta LaRC 104 0.193 0.273 0.019 0.057 

GG GRC 160 0.214 0.300 0 0.028 

HH LaRC 170 0.119 0.146 0.065 0.065 

II MSFC 210 0.146 0.229 0.033 0.087 

JJ GRC 220 0.197 0.236 0.021 0.066 

KK LaRC 245 0.122 0.159 0.047 0.070 

LL MSFC 255 0.095 0.154 0.062 0.062 

MM LaRC 280 0.132 0.184 0.023 0.039 

NN GRC 290 0.190 0.262 0 0.020 

OO MSFC 330 0.165 0.249 0 0.060 

PP LaRC 340 0.154 0.196 0.031 0.045 

QQ GRC 350 0.066 0.083 0.099 0.099 

RR remnanta GRC 356 N/A N/A N/A N/A 

a

Remnants are less than 10°.

TABLE VII.—ELECTRON SPECTROSCOPY FOR CHEMICAL ANALYSIS 

DATA FROM RELIEF RADIUS CRACK NEAR BOLT HOLE L OF S/N 120 

 Low oxide line (Nb3d5/2) High oxide line (Nb3d3/2) 

Sample 

location 

Position, 

eV 

Percentage Position, 

eV 

Percentage 

A 204.8 65.4 210.5 34.6 

B 203.9 71.9 209.5 28.1 

C a202.0 a83.7 210.3 16.2 

D 204.4 81.7 210.4 18.3 

E 204.4 71.8 210.5 28.2 

F 204.4 65.2 210.3 34.7 

a

Line position indicates metallic Nb. 
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77

Reaction Control System Thruster Cracking Consultation: NASA

Engineering and Safety Center (NESC) Materials Super Problem

Resolution Team (SPRT) Findings

Rebecca A. MacKay, Stephen W. Smith, Sandeep R. Shah, and Robert S. Piascik

Metals; Alloys; Niobium; Hydrogen embrittlement; Refractory metals; Testing and

evaluation; Stress corrosion cracking; Corrosion; Thruster; Fractography

Unclassified -Unlimited

Subject Category: 26

Rebecca A. MacKay, NASA Glenn Research Center; Stephen W. Smith, NASA Langley Research Center;

Sandeep R. Shah, NASA Marshall Space Flight Center; and Robert S. Piascik, NASA Engineering and Safety

Center, NASA Langley Research Center. Responsible person, Rebecca A. MacKay, organization code RX,

216–433–3269.

The shuttle orbiter’s reaction control system (RCS) primary thruster serial number 120 was found to contain cracks in

the counter bores and relief radius after a chamber repair and rejuvenation was performed in April 2004. Relief radius

cracking had been observed in the 1970s and 1980s in seven thrusters prior to flight; however, counter bore cracking

had never been seen previously in RCS thrusters. Members of the Materials Super Problem Resolution Team (SPRT)

of the NASA Engineering and Safety Center (NESC) conducted a detailed review of the relevant literature and of the

documentation from the previous RCS thruster failure analyses. It was concluded that the previous failure analyses

lacked sufficient documentation to support the conclusions that stress corrosion cracking or hot-salt cracking was the

root cause of the thruster cracking and lacked reliable inspection controls to prevent cracked thrusters from entering the

fleet. The NESC team identified and performed new materials characterization and mechanical tests. It was determined

that the thruster intergranular cracking was due to hydrogen embrittlement and that the cracking was produced during

manufacturing as a result of processing the thrusters with fluoride-containing acids. Testing and characterization

demonstrated that appreciable environmental crack propagation does not occur after manufacturing.